CORROSION BEHAVIOUR OF ALUMINISED STEEL AND CONVENTIONAL ALLOYS IN SIMULATED ALUMINIUM SMELTING CELL ENVIRONMENTS
by
Nan Xu
A Thesis Submitted for the Degree of Doctor of Philosophy
School of Materials Science and Engineering The University of New South Wales Australia
April 2002
CERTIFICATE OF ORIGINALITY
I hereby declare that this submission is my own work and that, to the best of my knowledge it contains no materials previously published or written by another person, nor material which to a substantial extent has been accepted for the award of any other degree or diploma at UNSW or any other educational institution, except where due acknowledgement is made in the thesis. Any contribution made to the research by others, with whom I have worked at UNSW or elsewhere, is explicitly acknowledged in the thesis.
I also declare that the intellectual content of this thesis is the product of my own work, except to the extent that assistance from other in the project’s design and conception or style, presentation and linguistic expression is acknowledged.
Nan Xu
DEDICATION
I dedicate this work to my mother RenZhuang Bai.
i
ACKNOWLEDGEMENTS
I wish to express my thanks to the following people for their assistance with this research project:
To Professor David J. Young for his valuable technical advice, idea and guidance provided throughout the whole project. I would also like to thanks Professor Young for giving me the opportunity to participate in this project and also for organising the funding for this project.
To all technical staff at the School of Materials Science and Engineering for their assistance.
To Comalco Pty. Ltd. for their technical support related to this project and supplying required materials.
To all other fellow members of the High-Temperature Materials Group for their assistance and helpful general discussion.
To my parents and my beloved wife for their encouragement and understanding, and support throughout my studies.
ii ABSTRACT
Aluminium smelting is a high temperature electrometallurgical process, which suffers considerable inefficiencies in power utilization and equipment maintenance.
Aluminium smelting cell works in the extreme environments that contain extraordinarily aggressive gases, such as HF, CO and SO2. Mild steel used as a structural material in the aluminium industry, can be catastrophically corroded or oxidized in these conditions. This project was mainly concerned with extending the lifetime of metal structures installed immediately above the aluminium smelting cells.
An aluminium-rich coating was developed on low carbon steel A06 using pack cementation technique. Yttria (Y2O3) was also used to improve the corrosion resistance of coating. Kinetics of the coating formation were studied. XRD, FESEM and FIB were employed to investigate the phase constitution and the surface morphology.
Together with other potentially competitive materials, aluminium-rich coating was evaluated in simulated plant environments. Results from the long time (up to 2500h) isothermal oxidation of materials at high temperature (800°C) in air showed that the oxidation resistance of coated A06 is close to that of stainless steel 304 and even better than SS304 in cyclic oxidation tests. Coated A06 was also found to have the best sulfidation resistance among the materials tested in the gas mixture contains SO2 at
800°C. Related kinetics and mechanisms were also studied. The superior corrosion resistance of the coated A06 is attributed to the slow growing a-Al2O3 formed. Low temperature corrosion tests were undertaken in the gas mixtures containing air, H2O,
HCl and SO2 at 400°C. Together with SS304 and 253MA, coated A06 showed excellent corrosion resistance in all the conditions. The ranking of the top three materials for corrosion resistance is: 253MA, coated A06 and SS304. It is believed that
iii aluminised A06 is an ideal and economical replacement material in the severe corrosive aluminium smelting cell environment.
iv TABLE OF CONTENTS
Dedication i
Acknowledgements ii
Abstract iii
Table of Contents v
List of Figures ix
List of Tables xvi
CHAPTER 1: INTRODUTION 1 1.1 INTRODUCTION 1 1.2 FAILURE ANALYSIS 3 1.3 CONCLUSION FROM FAILURE ANANLYSIS 11
CHAPTER 2: LITERATURE REVIEW 12 2.1 OXIDATION OF METALS AND ALLOYS 12 2.1.1 Introduction 12 2.1.2 Oxidation of Metals 12 2.1.2.1 Thermodynamic aspects 12 2.1.2.2 Mechanism and kinetics of scale formation 13 2.1.2.2.1 Parabolic law 13 2.1.2.2.2 Linear rate law 19 2.1.3 Oxidation of alloys 19 2.1.3.1 External scale formation 20 2.1.3.2 Influence of alloying elements 21 2.1.3.3 Internal oxidation 23 2.1.4 Oxidation in mixed gas 24 2.2 SULFIDATON OF METALS AND ALLOYS 27 2.2.1 Introduction 27 2.2.2 Properties of metal sulfides 27
v 2.2.3 Sulfidation of metals 28 2.2.3.1 Sulfidation of Iron 28 2.2.4 Effects of the sulfur pressure, temperature and gas 31 2.2.5 Sulfidation of binary alloys 34 2.2.5.1 Sulfidation of Iron-Chromium 34 2.2.6 Sulfidation of commercial alloys 38 2.2.6.1 Sulfidation of carbon steel 38 2.2.6.2 Sulfidation of low alloy steels 39 2.2.6.3 Sulfidation of ferritic stainless steels 39 2.2.6.4 Sulfidation of Austenitic stainless steels 40 2.3 HIGH TEMPERATURE CORROSION OF METALS AND ALLOYS IN
ATMOSPHERES CONTAINING HCL/CL2 43 2.3.1 Introduction 43 2.3.2 Thermodynamic aspects 43
2.3.3 High temperature corrosion of iron in atmospheres containing HCl/Cl2 44 2.3.4 High temperature corrosion of chromium in atmospheres containing
HCl/Cl2 45 2.3.5 The chlorination of alloys containing chroumium 46 2.3.6 Chlorination of alloys containing aluminium 49 2.3.7 Active oxidation mechanism 52 2.4 PACK CEMENTATION ALUMINIZIED COATING 54 2.4.1 Introduction 54 2.4.2 The pack cementation process 54 2.4.3 Procedure 55 2.4.4 Mechanism of coating formation 56 2.4.5 Process variables in pack cementation 60 2.4.6 Pack aluminizing on Ni-based and Fe-based alloys 60 2.4.6.1 Pack aluminizing on Ni-based alloys 60 2.4.6.2 Pack aluminizing of iron and commercial steels 69 2.4.6.3 High temperature corrosion resistance of iron aluminide 72 2.4.6.3.1 High Temperature Oxidation of Iron Aluminides 72 2.4.6.3.2 High Temperature Sulfidation and Oxidation/Sulfidation of iron aluminides 76
vi 2.4.6.3.3 High temperature chlorination of iron aluminide alloys 86
CHAPTER 3: EXPERIMENTAL PROCEDURES 93 3.1 SPECIMEN PREPARATION 93 3.2 PACK ALUMINIZING PROCESS 95 3.2.1 Pack Materials 95 3.2.2 Pack Preparation and Coating Deposition 96 3.2.2 Coating Analyses 97 3.3 HIGH TEMPERATURE OXIDATION 98 3.3.1 High Temperature Isothermal Oxidation 98 3.3.2 High Temperature Cyclic Oxidation 101 3.4 CORROSION TESTS IN MIXED GASES 101 3.4.1 Corrosion Tests in Gas Mixtures at Low Temperature 101 3.4.2 Corrosion Tests in Gas Mixture at High Temperature 102
CHAPTER 4: PACK ALUMINIZING COATING FORMED ON A1006 STEEL 104 4.1 INTRODUCTION 104 4.2 THERMODYNAMIC CALCULATION 104 4.3 RESULTS 105 4.3.1 Coatings Formed on Steel Using Different Activators without Yttria Deposition 105 4.3.2 Coating formed using NaF activator on steel pre-coated with yttria slurry 110 4.3.3 Surface morphologies of coating grown using NaF activator 114 4.3.4 Phase constitution of coatings formed for various deposition times 114 4.3.5 Coating growth kinetics 123 4.4 DISCUSSION 124 4.4.1 Effect of activator type on coating formed on A1006 steel 124 4.4.2 Surface morphology and kinetics of coating formation 125 4.4.3 Effect of yttria slurry deposit on coating constitution 127 4.5 CONCLUSIONS 128
CHAPTER 5: HIGH TEMPERATURE OXIDATION 129 5.1 INTRODUCTION 129
vii 5.2 RESULTS 129 5.2.1 Isothermal Oxidation 129 5.2.1.1 TGA Experiment 129 5.2.1.2 Long Time Oxidation Experiments 140 5.2.2 HIGH TEMPERATURE CYCLIC OXIDATION 146 5.3 DISCUSSION 153 5.3.1 Oxidation Kinetics for Materials Tested Using TGA 153 5.3.2 Scale morphology study 154 5.4 CONCLUSION 158 CHAPTER 6: HIGH TEMPERATURE SULFIDATION 159 6.1 INTRODUCTION 159 6.2 THERMODYNAMIC ASPECT 159 6.3 RESULTS 162 6.4 DISCUSSION 191 6.5 CONCLUSION 197
CHAPTER 7: LOW TEMPERATUER TESTS 198 7.1 INTRODUCTION 198 7.2 THERMODYNAMIC ASPECTS 198 7.3 RESULTS 203 7.4 DISCUSSION 228 7.5 CONCLUSION 235
CHAPTER 8: SUMMARY AND CONCLUSIONS 236
CHAPTER 9: REFERENCES 240
APPENDIX 1: LIST OF DATA USED IN COMPUTER CALCULATIONS 258
APPENDIX 2: XRD RESULTS FOR LONG TERM ISOTHERMAL OXIDATION EXPERIMENTS 260
viii LIST OF FIGURES
FIGURE PAGE
1.1 (a) Aluminium reduction cell-ferrous component degradation; 2
(b) Schematic diagram showing the high temperature
oxidation/corrosion occurring on the anode stub at the site
above the carbon anode of a reduction cell [1] and
(c) Corroded stub.
1.2 Schematic drawing of hopper and section where samples were 5
taken from.
1.3 Corroded samples sectioned from hopper. 6
1.4 (a) Optical image of the cross-section of sample 1 (etched by 8
5wt% nital);
(b) Optical image of the cross-section of sample 2 (etched by
5wt% nital);
(c) Optical image of the cross-section of sample 3.
1.5 (a) Thinned plate (sample 2) and scale at its HAZ (etched by 10
5wt% nital);
(b) Open surface of thinned plate (sample 2);
(c) Thicker plate (sample 2) and its HAZ (etched by 5wt% nital).
2.1 Ellingham diagram-Standard Gibbs free energy of formation of 14
selected oxides as a function of temperature [6].
ix 2.2 Schematic diagram of main phenomena and part-processes taking 15
place in the reaction of metals with single oxidant, e.g., oxygen
[7].
2.3 Simplified model for diffusion controlled oxidation [8]. 15
2.4 Schematic diagram of the formation of the external scale on the 21
binary alloy AB (B is the less noble element).
2.5 Hypothetical thermodynamic stability diagram of metal M in bi- 25
oxidant X2+Y2 environment.
2.6 Ellingham diagram for selected metal sulfides [58-59]. 2.9
2.7 Cross-section of scale formed on pure iron sulfidised at 973K for 30
22.5mins [65].
2.8 Sulfur partial pressure dependence of the parabolic rational rate 33
constant for the sulfidation of pure iron [67].
2.9 Sulfidation rate constant for Fe-Cr alloys at P =1atm and 37 S2
T=173K [82].
2.10 Cross-section of scale formed on Fe-14Cr alloy reacted with 38
P at 1173K [75]. S2
2.11 Sulfidation rates of chromium steels in H2S/H2S at 12 atm at 41
908K [91].
2.12 Scale formed on type 310 steel at P =1x10-7 atm and 1145K 42 S2
[97].
2.13 Diagram of change of weight losses of Cr2O3 exposed to reaction 47
gas for 2h [103].
x 2.14 Schematic overview of several reactions considering the induced 53
corrosion attack by the presence of HCl or Cl2 [129].
2.15 Schematic diagram of a pack cementation retort [140]. 57
2.16 Fe-Al binary alloy phase diagram [146]. 59
2.17 Variation of Ni-specimen surface composition and weight gain 62
with time in 4 wt% aluminium packs at 1000°C [134].
2.18 (a) Activator “circulation” and 67
(b) “condensation” models for Al transport [148].
2.19 Mixed mechanism of Al transport in the presence of both 68
activator-only, and activator and Al depleted zones [149].
2.20 Corrosion Kinetics for Fe-Al alloys: (a): Fe-5Al; (b) Fe-10Al; (c) 83
Fe-18Al; (d) Fe-28Al and (e) Fe-40Al [188].
2.21 Weight changes for several iron aluminide coatings and uncoated 88
austenitic alloys exposed in gas mixtures containing H2S with and
without HCl [195].
2.22 Weight Change data of Fe3Al-based alloys tested in a 1vol% HCl- 91
N2 gas mixture at 650°C [197].
3.1 Cross-section image of a specimen with pre-deposited Y2O3 95
slurry.
3.2 Schematic drawing of pack aluminizing process equipment. 97
3.3 Schematic drawing of thermogravimetric analyser 100
3.4 Schematic drawing of the high temperature apparatus 100
3.5 Schematic drawing of cyclic oxidation apparatus 102
xi 3.6 Schematic drawing of the apparatus used for low temperature 103
corrosion tests.
4.1 (a) SEM image of the cross-section of coating formed using 108
NH4Cl activator;
(b) Optical image of the cross-section of coating formed using
NaF activator (etched by 5% Nital).
4.2 (a) XRD scan result for coating formed using NH4Cl activator; 109
(b) XRD scan result for coating formed using NaF activator.
4.3 Aluminium concentration of coating formed using NH4Cl 110
activator at 900°C after 8hr.
4.4 EPMA result for coating formed using NaF activator after 8hr at 111
900°C.
4.5 Cross-sectional optical image of the of coating (etched by 1 part 112
of HF + 2 parts of Nital + 1 part of H2O) formed on A1006 steel
(pre-coated with yttria slurry) at 900°C after 8hr.
4.6 EPMA results for coating formed on A1006 steel (yttria slurry 113
pre-coated) at 900°C after 8hr.
4.7 Surface morphology of coating formed after different deposition 115
time.
4.8 Optical images of the cross-sections of coatings (etched) formed 119
for various deposition times at 900°C: (a): 0 hr, (b): 2hr, (c) 4hr,
(d) 6 hr and (e) 8hr.
4.9 XRD results of coatings formed for various deposition times. 120
xii 4.10 EPMA results for coatings formed for various deposition times: : 123
(a): 0 hr, (b): 2hr, (c) 4hr, (d) 6 hr and (e) 8hr.
4.11 (a) Diagram of coating thickness vs. time (using NaF activator) at 126
900°C ;
(b) Diagram of coating thickness2 vs. time (using NaF activator) at
900°C .
5.1 (a) Oxidation kinetics of A06, 2.25Cr1Mo, and SG cast iron 134
exposed to air at 800°C;
(b) Oxidation kinetics of 3Cr12 exposed to air at 800°C in TGA;
(c) Oxidation kinetics of SS304 and 253MAat 800°C in air;
(d) Oxidation kinetics of coated A1006 steel at 800°C in air.
5.2 Cross-sectional images of A1006 steel, 2.25Cr1Mo steel and SG 137
cast iron after oxidized (TGA) at 800°C.
5.3 SEM images for 3Cr12, SS304 and 253MA oxidized in air at 139
800°C in TGA experiments.
5.4 Weight changes of coated A1006, SS304 and 253MA steel 141
oxidized in air at 800°C for different exposure times.
5.5 Oxidation kinetics of coated A1006 and SS304 steel in air at 142
800°C .
5.6 SEM images for coated A1006 steel oxidized in air at 800°C for 145
2500 hours.
5.7 SEM images for coated A1006 steel oxidized in air at 800°C for 147
2500 hours.
xiii 5.8 SEM images for SS304 (a) and 253MA (b, c) oxidized in air at 148
800°C for 2500 hours.
5.9 Weight changes of bare A1006 steel, SS304, 253MA and coated 150
A1006 steel in cyclic oxidation experiments.
5.10 Cross-sectional images for bare A1006 steel, 253MA, SS304 and 152
coated A1006 steel oxidized in cyclic oxidation experiments.
5.11 Diagram of weigh gain^2 vs. time for (a) A1006 steel and (b) 155
2.25Cr1Mo steel.
5.12 Oxidation kinetics for SG cast iron in the initial stage. 156
6.1 Stability diagram of M-S-O system at 800°C 161
6.2 Weight gain of samples in sulfidation experiments at 800°C. 163
6.3 Kinetics of materials in sulfidation experiments at 800°C. 165
6.4 Cross-sectional images for A1006 steel in high temperature 171
sulfidation experiments at 800°C after different exposure time.
6.5 Cross-sectional images for 2.25Cr1Mo steel in high temperature 174
sulfidation experiments at 800°C after different exposure time
6.6 Cross-sectional images for SG cast iron in high temperature 178
sulfidation experiments at 800°C after different exposure time
6.7 Cross-sectional images for 3Cr12 steel in high temperature 180
sulfidation experiments at 800°C after different exposure time
6.8 Cross-sectional images for SS304 steel in high temperature 183
sulfidation experiments at 800°C after different exposure time
6.9 Cross-sectional images for 253MA steel in high temperature 186
sulfidation experiments at 800°C after different exposure time.
xiv 6.10 Cross-sectional images for coated A1006 in high temperature 189
sulfidation experiments at 800°C after different exposure time.
6.11 SEM images for materials sulfidized at 800°C for 16 hours. 190
6.12 SEM image for 253MA sulfidized at 800°C for 16 hours. 191
6.13 Isothermal section for the Fe-Cr-O phase diagram at 1200°C [8]. 195
7.1 Stability diagrams for materials tested in M-Cl-O and M-S-O 202
systems at 400°C.
7.2 Kinetics for materials tested at 400°C in different gas mixtures. 209
7.3 Weight changes for materials exposed to different gas mixtures 214
for 500h at 400°C.
7.4 Optical images for some materials exposed to 4 different gas 224
mixtures at 400°C for 500 hours.
7.5 SEM images for materials exposed to gas 3 for 500 hours at 227
400°C .
7.6 Schematic diffusion path for formation of a scale consisting of 230
chloride and oxide underneath single-phase oxide.
7.7 Schematic drawing of the scale structure of pure copper and Cu-S 233
alloy formed in gas 3 and gas 4 at 400°C.
A.1 XRD results for SS304, 253MA and coated A1006 steel oxidized 268
in air at 800°C after different exposure times.
xv LIST OF TABLES
TABLE PAGE
1.1 XRD results for corrosion products collected from hopper 4
skirting.
1.2 EDAX results for corrosion products and metals. 9
2.1 Gas compositions used in the study of Bakker [122]. 50
2.2 Alloy compositions [122]. 51
2.3 Aluminium activities in Al-Fe system at 900°C [145]. 58
2.4 Equilibrium partial pressure of gases in packs activated by AlF 3, 66
NaF, NaCl and NH4Cl at 1000°C [135].
2.5 Composition, Temperature range and Oxygen partial pressure of 81
the test gases [187].
2.6 Parabolic Rate constants (g2 cm-4 sec-1) and Apparent Activation 84
Energies of Fe-Al Alloys [188].
2.7 Alloy compositions in the study of Saunders et al. [193]. 88
2.8 Calculated partial pressures (Total pressure = 101325 Pa) [193]. 89
2.9 Experimental conditions [196]. 91
2.10 Alloy composition (wt%) tested in simulated gasifier gases 92
[196].
3.1 Table of the compositions of materials investigated (wt%). 94
3.2 Proportions of pack constituents used in pack aluminizing 96
process.
3.3 Experimental conditions for mixed gas corrosion. 103
xvi 4.1 Partial pressures of species and condensed phases in packs using 107
different activators.
4.2 XRD results for coating formed on A1006 steel (pre-coated with 116
yttria slurry) at 900°C after 8hr.
5.1 XRD results for materials oxidized in air at 800°C in TGA 135
experiments
5.2 Parabolic rate constants for oxidation in air at 800°C 140
5.3 XRD results for oxide scales grown in air at 800°C 143
5.4 XRD results for materials oxidized in cyclic oxidation 149
experiments
6.1 Equilibrium partial pressure and activities for species in high 160
temperature sulfidation experiments at 800°C.
6.2 Parabolic sulfidation rates at 800°C. 165
6.3 XRD results of in situ scales grown in high temperature 166
sulfidation experiments.
7.1 Equilibrium partial pressure and activities for species in high 199
temperature sulfidation experiments at 400°C.
7.2 XRD results of materials reacted in different gas mixtures at 211
400°C.
7.3 Vapour pressures of metal chlorides at 400°C. 231
7.4 Calculated equilibrium Cl 2 potential at layer boundaries in 232
copper oxide scale at 400°C .
A.1 Table of the Gibbs Free Energy used in NH4Cl pack. 258
A.2 Table of the Gibbs Free Energy used in NaF pack. 259
xvii Introduction
CHAPTER ONE
Introduction
1.1 INTRODUCTION
Australia is the world largest bauxite and alumina producer, and the fourth largest aluminium producer. The export value of alumina and aluminium is high, and its continued success depends critically on both price and quality due to the intensive competition.
Aluminium smelting is a high temperature electrometallurgical process, which suffers considerable inefficiencies in power utilization and equipment maintenance.
The cost results from the design of the smelting cell and its extreme operating conditions. Figure 1.1 shows a schematic diagram of the aluminium-smelting cell. The gas temperature in the cell ranges from 200°C to 800°C in the vicinity of the anode stubs and down to from 50°C to 340°C near the deck plate. The atmospheres contain extraordinarily aggressive gases, such as HF, CO and SO2 [1]. In this kind of environment, mild steel, which is widely used as a structural material in aluminium smelters, can be catastrophically corroded. For example, an anode stub used for supporting the carbon anode and for conducting electrical current was found to be corroded after a relatively short time [1]. The corroded section of the stub is shown in
Figure 1b [1] and Figure 1c. The deck plate around the rim of the pot and the alumina hopper are other sections, which can also be corroded severely. The corrosion products can be introduced into the cell bath, and then contaminate the purity of the aluminium products. This has a substantial effect on the aluminium industry profitability, since the
1 Introduction
Alumina hopper Gussets Fume deflectors Stub Deck-plates
S ide wall Liquid aluminium pad
(a)
Anode stub Corrosion site Crust
Molten bath
(b) (c)
Figure 1.1 (a) Aluminium reduction cell-ferrous component degradation;
(b) Schematic diagram showing the high temperature oxidation/corrosion
occurring on the anode stub at the site above the carbon anode of a
reduction cell [1];
(c) Corroded stub.
2 Introduction
aluminium quality is a key aspect of the export market. Therefore, finding a suitable material which can withstand these severe high temperature corrosion conditions is highly desirable in the aluminium smelting industry.
1.2 FAILURE ANALYSIS
To obtain a general understanding of the severe corrosion problems encountered, preliminary analysis on samples collected from an aluminium smelting cell were carried out.
Several samples were collected from a hopper which was corroded after a relatively short time. A schematic drawing and corresponding corrosion samples for the hopper are shown in Figure 1.2 and Figure 1.3, respectively.
It should be pointed out that the whole surface of the hopper outside was covered by a layer of white coloured powder. XRD results showed that the main component of the powder was alumina, and therefore it is concluded that the white coloured powder came from the raw materials. XRD results for corrosion products collected from the hopper skirting are shown in Table 1.1. Noting that Al2O3 and sodium salts are not corrosion products. It is concluded that the steel is not only oxidising but is also being corroded by sulfur-containing species. These XRD results are consistent with other researchers’ results, which explained the presence of iron sulfide and iron sulfate [1].
When the working temperature of the hopper is below 330°C, previous formed Fe2O3 scale can absorb the by-product – gaseous SO2 and react with O2 to form iron sulfate
(Eq. 1.1). Since the P of Fe (SO ) is as low as 10-15 atm at 200°C, Fe O can be SO 2 2 4 3 2 3 regarded as an excellent sorbent for SO2 at the low temperatures. However, Fe2(SO4)3 will decompose and release SO2 and O2 when the temperature is higher than 430°C.
Released SO2 can diffuse through the scale and reach the scale/steel interface where the
3 Introduction
partial pressure of oxygen and SO2 are low. Iron can be attacked by both sulfur and oxygen as shown in Eq. 1.2 and Eq. 1.3. Fe2O3 at the scale outer surface will act as a
SO2 sorbent once again when required temperature is reached.
Fe2O3 + 3SO2 +3/2O2 = Fe2(SO4)3 (1.1)
5Fe + 2SO2 = 2FeS + Fe3O4 (1.2)
4Fe + 2SO2 = FeS2 + Fe3O4 (1.3)
Severe local attack was observed, as shown in Figure 1.3. Although the original thickness of the steel sheet was 3mm, some parts of the hopper structure were completely penetrated (samples 1 and 3). As seen in Figure 1.3a, only the welding material was left in Sample 1.
Table 1.1 XRD results for corrosion products collected from hopper skirting
Sample ID Phases identified
Fe2O3
FeS2
Fe1-xS Hopper skirting Fe2(SO4)3
Al2O3
Na2S2O4
4 Introduction
irting sk
Hopper
Sample 3
Sample 2
Schematic drawing of hopper and section where samples were taken from
Figure 1.2 Sample 1
5 Introduction
(c)
Sample 3
(b)
Sample 2
Corroded samples sectioned from hopper
Figure 1.3
(a)
Sample 1
6 Introduction
Samples were then cross-sectioned and the results are shown in Figure 1.4. Welded samples were further etched with 5% Nital to reveal the grain structure, and results are shown in Figure 1.4. The T-joint (Figure 1.4a) is seen to be defective: full penetration had not been achieved on either side. The resulting cavity contains a deposit and its inner walls have been oxidised. The butt joint (Figure 1.4b) is sound, but extensive metal loss had occurred to one of the two plates. The thinned plate (Figure 1.5a) had been consumed uniformly over its surface, although corrosion product is evident mainly in the notch where the metal heat affected zone (HAZ) had been consumed (Figure
1.5b). It is likely that the scale has been retained by the notch, whereas scale has fallen off the open surface. Examination of the thicker plate (Figure 1.5c) shows no significant acceleration of attack in the HAZ relative to the open surface. In fact, thicker portions of scale are seen on the open surface. It is clear that the weld metal was more resistant to attack than the parent metal. However, no further investigation was carried out since this was not the purpose of this project. For Sample 3 (Figure 1.4c), its tip section indicates that the plate had been completely penetrated. Porous, multiple- layered scale was found to be mainly iron oxides together with a trace amount of sulfur at the scale/steel interface.
Samples were analysed using EDAX and results are shown in Table 1.2. It should be pointed out that EDAX technique is not sensitive to oxygen. Furthermore, fluoride could not be detected in this case because of interference by one of the Fe lines with the principal F line. Further chemical analysis for fluoride indicated a significant amount of fluoride (4.19 wt %).
It is concluded that the corrosion products are principally iron oxides, but that significant sulfur corrosion had also occurred. In particular, sulfur was located at the scale/metal interface, indicating that sulfur species had penetrated the non-protective
7 Introduction
(a)
(b)
(c)
Figure 1.4 (a) Optical image of the cross-section of sample 1 (etched by 5wt% nital);
(b) Optical image of the cross-section of sample 2 (etched by 5wt% nital);
(c) Optical image of the cross-section of sample 3.
8 Introduction
scale. Under these circumstances, iron oxide and sulfide will grow simultaneously, a process known to be much faster than simple oxidation of iron.
Table 1.2 EDAX results for corrosion products and metals
Sample ID Area Elements detected
Deposit inside weld defect Ti, Si, Mg, Ca, K, Mn, Fe
Scale inside weld defect Fe
1 Innermost scale on thinned Fe, S (trace)
Layered scale on plate Fe
Weld bead Fe, Mn, Si
Parent metal Fe
Innermost scale at weld/parent metal Fe, S (trace)
junction
2 Layered scale Fe
Weld bead Fe, Mn, Si
Parent metal Fe
3 Scale Fe
9 Introduction
(a)
(b)
(c)
Figure 1.5 (a) Thinned plate (sample 2) and scale at its HAZ (etched by 5wt% nital);
(b) Open surface of thinned plate (sample 2);
(c) Thicker plate (sample 2) and its HAZ (etched by 5wt% nital).
10 Introduction
1.3 CONCLUSION FROM FAILURE ANANLYSIS
Although the corrosion behaviour of pure metals and alloys at high temperature has been investigated for many years [2, 3], the mechanisms and kinetics of the relevant reactions under the operating conditions of an aluminium-smelting cell have not been well explained and understood.
The present project was mainly concerned with extending the lifetime of metal structures installed immediately above aluminium smelting cells. Successful alloys used in this situation must form an oxide scale which is stable with respect to significant levels of HF and SO2. The only practical oxides are Cr2O3 and Al2O3. It is known [4] that Cr2O3 scale is permeable to a number of secondary oxidants. On the other hand,
Al2O3 scale has superior resistance to chlorine or sulfur containing atmospheres and is known to be more resistant to sulfur than Cr2O3. Accordingly, the present project was concentrated chiefly on developing an aluminium-rich coating on mild steel for service on hardware components above smelting pots. Together with other potentially competitive materials, the aluminized coating was evaluated in simulated plant conditions, reaction mechanisms investigated and a ranking of the candidate materials arrived at.
11 Literature Review
CHAPTER TWO
Literature Review
2.1 OXIDATION OF METALS AND ALLOYS
2.1.1 Introduction
The high temperature oxidation of metals and alloys is a metal consuming process.
The rate of oxide formation depends on the reaction temperature and the reactive gas partial pressure in the environment. In this chapter the mechanism and kinetics of metal oxidation are reviewed.
2.1.2 Oxidation of Metals
2.1.2.1 Thermodynamic aspects
When a metal M is exposed to an oxidant, such as O2, at high temperature, the metal oxide will be formed via the following reaction:
y (2.1) xM s)( O g)( =+ M O 2 2 yx
According to the second law of thermodynamics, the driving force of this reaction can be written as [5]:
a (2.2) M Oyx D = DGG ° + RT ln( y ) a x P 2 M O2
Where DG and DG°are the Gibbs free energy change and the standard free energy of formation, respectively, R is the universal gas constant, T is the absolute temperature,
12 Literature Review
a and a are the activity of metal oxide and metal respectively and P is the M xO y M o2 partial pressure of oxygen. When DG < 0, spontaneous reaction is expected.
The activities of M and MxOy can be taken as unity in most oxidation reactions, and equation (2.2) can be used to express the dissociation partial pressure of oxide at equilibrium (DG = 0):
2DG° (2.3) PO = exp( ) 2 yRT
If the oxidant partial pressure in the environment is higher than the dissociation partial pressure of the oxide, then oxidation of metal M will occur. In the Ellingham diagram (Figure 2.1) [6], straight lines are drawn to show the standard free energy of formation of selected oxides as functions of temperatures. In the diagram, the lower the reaction line is, the more stable the oxide is, e.g. Cr2O3 is less stable than Al2O3, but is more stable than Fe2O3.
2.1.2.2 Mechanism and kinetics of scale formation
2.1.2.2.1 Parabolic law
The scale formation process for metal M exposed to an oxidant like O2, includes serval stages such as: the adsorption of oxygen molecules on the metal, dissociation of
O2 into atoms and the formation of an oxide MxOy. The initial oxide nuclei will form at specific locations on the metal surface. Continuous oxide layer will be formed by their subsequently growth. This process is schematically illustrated in Figure 2.2 below [7].
When reactants are separated by a scale without physical discontinuities, such as cracks and pores, the diffusion of one or both reactants though the lattice and along grain boundaries will occur to support the further reaction. Both the chemical and electrical
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Figure 2.1 Ellingham diagram-Standard Gibbs free energy of formation of
selected oxides as a function of temperature [6].
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Figure 2.2 Schematic diagram of main phenomena and part-processes taking
place in the reaction of metals with single oxidant, e.g., oxygen [7].
Figure 2.3 Simplified model for diffusion controlled oxidation [8].
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potential gradients provide the driving force for solid-state diffusion. If the ionic transport across the developing oxide layer controls the scaling rate, and thermodynamic
equilibrium is established at all interfaces, then the outward cation flux, j M 2 + is equal and opposite to the inward cation vacancy flux, j . Figure 2.3 [8] shows a simplified v m model for diffusion-controlled oxidation. To a first approximation,
-CC ''' M VV M 2 + jj =-= D (2.4) M VM VM X where X is the thickness of the oxide, D is the diffusion coefficient for cation VM vacancies, and C" and C' are the vacancy concentrations at the scale-gas and scale- VM VM metal interfaces respectively.
In view of the thermodynamic equilibrium established at each interface, CC''- ' VVM M is constant, thus equation (2.4) can be used to obtain:
''' dX -CC VV VJ == VD M M (2.5) dt Vm VM X
Where V is the volume of oxide formed per unit of flux, and t denotes time. Defining
CC''- ' k¢= D VVM M , we obtain: VM const.
dX k¢ (2.6) = dt X which upon integration yields,
2 X =kp¢t + C (2.7)
¢ ¢ ’ 2 -1 where k p is the parabolic rate constant (i.e. k p=2k with units cm .sec ), and C is a constant.
Equation (2.7) can also be expressed in terms of weight gain:
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2 (DW/A) =kpt + C (2.8)
where DW/A is the oxidation weight gain per unit surface area, and kp is the parabolic rate constant with units g2.cm-4×sec-1. The above equations apply to the situation when reactant diffusion through the oxide is the rate-controlling step. The parabolic rate constant for iron oxidation at 1000°C and 1 atm is 4.8´10-7 g2 cm-4 sec-1 [9].
The oxidation of iron in air or oxygen at low t emperatures has been studied by many researchers [10-15]. Different kinetics for iron oxidation were suggested below about 673K, where a two-stage oxidation was observed by thermogravimetry [16, 17].
Several rates laws were used to describe the oxidation behaviour in the first stage, such as logarithmic [18-21], double-logarithmic [22, 23] and parabolic. However, the oxidation in the second stage has not been fully understood. The iron oxidation rate was also found to be affected by several factors: oxygen partial pressure and heat treatment. Because of the slow oxidation rate of iron below about 673K, Rutherford backscattering spectroscopy (RBS) was employed by Sakai et al [24, 25] to study the oxidation of iron in the second stage between 523 and 673K in air, by measuring the thickness of the scale formed. At 673K, a single-stage oxidation was found, whereas below 648K, the two-stage reaction occurred. The scale thickness formed at 648K was thicker than that formed 673K after 1 hour exposure. This was explained as being due to the simultaneous growth of magnetite and hematite from the very beginning of the reaction at 673K. Below 648K, magnetite grows rapidly in the first stage until hematite covered it. The growth of magnetite at 648K was faster than the simultaneous growth of magnetite and hematite because the formation of dense hematite slowed down the total oxidation rate.
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High temperature oxidation of copper is a well studied subject [26-29]. The kinetics of the high temperature oxidation of copper are parabolic and depend on the temperature and the ambient partial pressure of oxygen. The oxide scale consists of either Cu2O or Cu2O+CuO. Since the diffusion in CuO is much slower than that in
Cu2O, Cu2O is the main part of the two-layered scale.
The theory of oxidation, which makes calculation of the parabolic rate constant possible, was developed by Wagner initially [30]. The following equation for kr is related to the defect properties of the oxide directly:
P" 1 O 2 Zc kr = ' ()DDM + o d ln PO (2.9) òP 2 2 O2 Za
2 -1 where kr is the parabolic rational rate constant (cm sec ), Zc and Za are the cation and
2 -1 anion charges respectively, DM and Do are the self diffusion coefficients (cm sec ) of the metal and oxygen respectively, P" and P' are evaluated at the external and O2 O2 internal scale surfaces, respectively. The relationship between kr and kp can be expressed by [8]:
2 éVZ ù 1 x (2.10) kr = ê ú k p 2 ë M x û
where V is the equivalent volume of scale, Zx is the valency of X, and Mx is the relative atomic mass of non-metal X, such as oxygen and sulfur, etc.
Wagner’s theory is based on the following assumptions [8]:
· The oxide layer is well adherent and compact;
· Diffusion of ions or electrons through the layer is the rate-controlling step;
· At both metal-scale and scale-gas interfaces, thermodynamic equilibrium is
established;
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· The oxide shows only negligible deviation from stoichiometry;
· Oxygen solubility in the metal can be neglected;
· The thickness of the scale is greater than the distances over which space charge
effects (i.e. electrical double layer) occur;
· Thermodynamic equilibrium is established locally throughout the scale.
The theory has been well proven by the experimental results measured for the oxidation of cobalt and sulfidation of iron.
2.1.2.2.2 Linear rate law
When a phase boundary process is the rate-determining step for the reaction, metal oxidation obeys the linear rate law.
X=klt (2.11)
-1 where kl is the linear rate constant with units cm sec . The linear rate law is commonly observed when oxidation is carried out in an atmosphere, which is diluted by, inactive or inert gas or if oxide formed is not protective. In the former case, the process of oxygen delivery to or adsorption onto the scale is so slow that it becomes the rate- controlling step rather than diffusion through the scale. In the latter case, a boundary reaction is the rate-controlling step rather than any mass transport process [8, 31].
2.1.3 Oxidation of alloys
Oxidation of alloys has been investigated by many researchers [32-36]. It has been shown that much of the scale growth description developed for the oxidation of pure metals can be applied to the oxidation of alloys, but the oxidation of alloys is, of course, much more complex. The alloying elements may [37, 38]:
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· have different affinities for oxygen caused by the different free energies of oxide
formation;
· cause the formation of ternary or higher oxides;
· exhibit a degree of solid solubility between the oxides;
· have different diffusivities in the alloy and different cationic mobilities in the
oxide phases;
· cause sub-surface precipitation of oxides of one or more alloying elements
(internal oxidation).
2.1.3.1 External scale formation
The various possible morphologies of scale formation on binary alloys have been reviewed by Bastow et al [39]. The effects of alloy compositions and reaction mechanisms on the composition and microstructure of the external scale were also examined.
Suppose a binary alloy AB (in which B is the solute and is less noble than A) is exposed to an oxidising environment at high temperature. It is possible to form different kinds of scale (Figure 2.4):
· an entirely AO scale on the alloy surface with the enrichment of B at the alloy-
scale interface caused by the reduced concentration of A in the alloy;
· a uniform BO scale on the alloy surface when the concentration of element B is
sufficiently high;
· a scale including both AO and BO since the following reaction occurs:
B+AO=BO+A (2.12)
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They can also form a ternary compound AxByO, solid solution (A, B)O, external and internal oxides depending on their properties and miscibility.
2.1.3.2 Influence of alloying elements
Elements such as chromium, nickel and molybdenum are added with various combinations into alloy steels to reach special requirements of strength, toughness and corrosion resistance. The influence of a few elements on the scaling behaviour of steel is reviewed:
AO (rich in B) BO (rich in A) AO+BO
A-B A-B A-B
(a) (b) (c)
Figure 2.4 Schematic diagram of the formation of the external scale on the binary
alloy AB (B is the less noble element).
Carbon: Carbon is the most common element in the steel and its effect on oxidation has been well studied [40-42]. Carbon reacts with iron oxide to produce CO gas at the scale/alloy interface. Whether it can decrease the reaction rate or not depends on if it can hinder the transport of the iron or oxygen ions or not. The origin of oxide nodules formed on Fe-3.4%C alloys at 760°C in low purity argon (containing100ppm oxygen) was studied by Merchant [43]. It was found that the oxide nodules formed during the very beginning stages of the oxidation in and around the cavities left by the denuded graphite inclusions. The new nodules appeared above the cavities, while the old ones grew in size along the length or width of the cavities, on further oxidation. The cavities
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were eventually completely covered with scale. Further growth took place on the surface of the sample until the nodules were connected with one another. It was believed that the nucleation of oxide nodules in the Fe-C base alloys was due to stress concentration near the edge of the cavities denuded of graphite. The oxidation kinetics of the Fe-C base alloy [44] were found to be a combination of the oxidation of iron and carbon in the initial stage. Although the oxidation of carbon increased with carbon content, it was also affected by the graphite morphology. If the graphite could be contacted by oxidant directly, its oxidation took place readily; however slower rates were observed when carbon was dissolved in the metallic matrix.
Chromium: Chromium increases hardenability and tensile strength. Because of the greater thermodynamic stability of Cr2O3 comparing with that of iron oxides, it can reduces the oxidation rate of steel when added at levels greater than 1.25 wt%. The formation of wustite can be prevented above 570°C when the chromium content in the steel is sufficiently high [45]. The reason for the reduced oxidation rate is regarded as the result of the formation and development of the Fe-Cr spinel (FeCr2O4) next to the base metal [46].
Aluminium: It is well known that the addition of aluminium to iron results in a significant increase in oxidation resistance [47], due to the formation of a uniform alumina layer. The nature of these layers appears to be different at different steel aluminium levels, temperature and oxidant activity [48].
Nickel: Nickel is more stable than iron and is rejected into the steel ahead of the iron oxide scale. It tends to concentrate at the steel/oxide interface because of its slow diffusion rate in the steel. When critical nickel level has been reached at the alloy surface for the particular environment (e.g. temperature, time and nickel content), the
22 Literature Review
formation of FeO becomes thermodynamically unfavourable, and Fe3O4 becomes the major oxide. As a consequence, the oxidation rate is reduced. It is thought that the low outwards diffusion rate of iron could not maintain growth of a FeO layer, so that its consumption in forming Fe3O4 occurred [49].
2.1.3.3 Internal oxidation
When oxygen diffuses into an alloy and sub-surface precipitation of the oxides of one or more alloying elements results, it is said that internal oxidation occurs. Internal oxidation has been reviewed by Wagner [50], Swisher [51] and Rapp [52]. The following criteria for internal oxidation have been proposed as the necessary conditions for the occurrence of the internal oxidation:
· The standard free energy change of the formation (per mole O2) for the solute
metal oxide must be more negative than that of the base metal;
· DG for the formation of the solute metal oxide must be negative, the oxidant
solubility and diffusivity in the base metal must be significant to establish the
necessary oxidant activity at the reaction front;
· The solute concentration of the alloy must be lower than that required for the
transition from internal to external oxidation;
· At the beginning of the oxidation, no surface layers may prevent the dissolution
of oxygen into the base alloy.
The internal oxidation reaction is initiated by the absorption of oxidant into either the external surface of the metal or at the alloy-scale interface (if an external scale is present). The oxidant then diffuses inward through the metal matrix and reacts with the solute alloy element to form oxide precipitates.
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The oxide front will move inward as the result of inward oxygen diffusion and the reactive alloying element outward diffusion. The rate-controlling step is oxygen diffusion into alloy and hence the kinetics of the internal oxidation are parabolic [35].
When the outward flux of solute to the alloy surface is enough to react with the oxidant and form a uniform oxide layer, the transition from internal oxidation to external oxidation occurs. This transition to external oxidation is the basis for Fe-and
Ni-base alloys which have sufficiently high concentration of a solute (such as Al, Cr or
Si) to form a stable and protective external oxide layer and hence prevent the oxidation of the metal matrix. The above process is known as selective oxidation [50]. Since the formation of the external or internal oxidation is a function of both the flux of inward diffusing oxygen and outward diffusing solute, the external scale formation could be increased either by increasing the outward diffusion of solute through grain boundaries and dislocation cores or by decreasing the flux of oxygen.
As a result of the external scale formation, the thickness of the internal oxidation zone would be less than that without external scale. Internal oxidation is usually, but not always undesirable. It would be expected that the surface scale might be more strongly keyed to the alloy by the sub-scale, and thereby be more protective due to its decreased susceptibility to spallation [53].
2.1.4 Oxidation in mixed gas
In a multi-oxidant environment, alloys may form a mixed primary oxidant phase, and even react with the second oxidant to form a different product. The effect of a second oxidant on the oxidation process can be explained as follows:
· the second oxidant may diffuse through the scale alloy interface and form
internal precipitates in the alloys;
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· the second oxidant may also modify scale transport properties by segregation to
the oxide grain boundaries, and;
· the second oxidant may enter the lattice of the oxide to from its own reaction
product, modifying the scale’s transport properties by acting as a short-circuit
path for cation transportation [54].
The oxidation product(s) formed on the alloy surface can be understood using a stability diagram, in which the regions of stability of the possible phases in equilibrium with the ambient gas mixture are shown. Figure 2.5 shows a schematic thermodynamic stability diagram of metal M in a bi-oxidant X2+Y2 environment.
MY
M MX
p Log ( X 2 )
Figure 2.5 Hypothetical thermodynamic stability diagram of metal M in bi-oxidant
X2+Y2 environment. where MX and MY are the corrosion products. The boundary between the MX and MY phase fields responds the following displacement-type reaction.
1 1 MX + Y ® MY + X (2.13) 2 2 2 2
Assuming unit activity in the solid phases M, MX and MY, the activity/partial pressure of the oxidant gases at the different phase boundaries can be written as follows:
25 Literature Review
1 DG ° a =P 2 = exp( MX ) (2.14) x x 2 RT
1 DG ° a =P 2 = exp( MY ) (2.15) y Y2 RT
1 2 a é PY ù DDGG- Y = ê 2 ú = exp( MY MX ) (2.16) ax Px RT ëê 2 ûú
The stability diagrams above can only be used at a starting point or guideline to predict the possible products. In practical, in addition to thermodynamic properties, the reaction kinetics can affect the reaction process and hence the scale products.
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2.2 SULFIDATION OF METALS AND ALLOYS
2.2.1 Introduction
The sulfidation of metals and alloys has been studied by a number of investigators because gaseous sulfur can lead to severe corrosion in a numbers of industries, such as oil and petrochemical. Sulfidation reactions exhibit the same kinetic phenomena as are observed for oxidation, but fewer studies have been made in the field of the basic kinetic phenomena involved in sulfidation of metals and alloys at high temperature [8].
Previous reviews of the sulfidation of metals and alloys have been made by
Kubaschewshi and Von Goldbeck [55] in 1954 and then in 1969 by Strafford [56]. A review of the sulfidation of iron and iron based alloys has been provided by Young [58].
In this section, the sulfidation of metals and alloys has been described using classical diffusion-controlled theory. The sulfidation of some commercial steels is also discussed.
2.2.2 Properties of metal sulfides
To study the thermodynamic properties of sulfides, the Ellingham diagram of several metal sulfides has been constructed by Shatynski [58], see Figure 2.6 (Data of the chromium sulfides from the work of Rau [59]). It is apparent that sulfides are less stable than oxides. The general physico-chemical properties of transition metal sulfides have been reviewed by Rao and Pisharody [60]. The non-stoichiometry in the sulfides of Fe, Co, Ni and Cr has been reviewed by Halstead [61]. Mrowec et al reviewed the defect structure and transport properties of selected sulfides and the sulfidation kinetics and mechanisms of some metals and alloys [62].
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Basically, there is a much larger degree of non-stoichiometry in sulfides than in the analogous oxides. Because of the larger anion size and lower lattice energy of the sulfides, it is easier for point defects to be created and hence deviation from stoichiometry arrived at. As the result of the higher density of lattice defects, a more poorly protective scale forms [63]. One good example is provided by chromium sulfides that possess a considerable range of non-stoichiometry [59, 64]. Consequently, the protection from sulfidation achieved by chromium additions is very limited. The non-stoichiometry or defect type of Al2S3 is not known. There is a little diffusion data for aluminium sulfide, chromium and aluminium-containing thio-spinels. It is also found that group VIII metals are quite unsuitable for high temperature service in sulfur- containing atmospheres, since most scales form low melting metal–metal sulfide eutectics.
2.2.3 Sulfidation of metals
2.2.3.1 Sulfidation of Iron
When pure iron is exposed to sulfur-containing atmospheres at temperatures below the Fe-S eutectic, a compact, tightly adherent scale is formed in a relative short term. If the value of P is sufficiently high, the scale consists of a thin surface layer of FeS S2 2 over a thick layer of Fe1-dS, and only the mono-sulfide phase is formed if the value of
P is low. A cross-section of scale formed on pure iron reacted with P =8´10-8atm S2 S2 for 22.5 mins at 973K is shown in Figure 2.7 [65].
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Figure 2.6 Ellingham diagram for selected metal sulfides [58-59].
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Figure2.7 Cross-section of scale formed on pure iron sulfidised at 973K for 22.5mins
[65].
It has been found that the growth of Fe1-dS follows parabolic kinetics, solid-state diffusion being the controlling-step. This reaction can also be described using the theory of Wagner [30]. Because the electronic conductivity characteristics of Fe1-dS are metallic in nature [66], and the self-diffusion coefficient of sulfur, DS is much less than that of iron, DFe, the flux of iron through a growing sulfide scale is predicted from
Wagner’s theory to be given by:
Z ln ad Fe S (2.17) Fe = cJ DFeFe 2 Z S dx
where J is the flux, c is the concentration, Z is the effective valence, aS is the activity of sulfur and x is the position coordinate within the scale normal to the reacting surface.
Integrating equation (2.17), one obtains:
" a S Z = ck Fe D ln ad (2.18) r eq òa ' Fe S S Z s
' " where ceq is the average concentration of metal in the scale, aS and aS are the sulfur activities at the metal-scale and scale-gas interfaces, respectively. The value of kp can then be found using Eq. (2.10). Fryt et al [67, 68] found the calculated values of kp to 30 Literature Review
Table 2.1 Gas compositions used in the study of Bakker [122].
1 2 3 4 5 6 6A 7 8
H2O 0 1 2 3 5 6.5 6.5 10 15
H2 32 30.2 29.2 28.2 28.2 28.2 43 28.2 30
CO 64 64 64 64 60 59 44.4 52 45
CO2 3.2 4 4 4 6 6 5 9 9.2
H2S 0.8 0.8 0.8 0.8 0.8 0.8 0.6 0.8 0.8
HCl 0, 0.04 0.04 0.04 0-0.0012 0.04 0 0.04 0.04 0, 0.04
Log PO2 -29 -28.4 -27.8 -27.4 -27 -26.8 n.d. -26.4 -26
Log PS2 -9.3 -9.3 -9.3 -9.3 -9.3 -9.3 n.d -9.3 -9.3
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be in very close agreement with the measured values of kp for 873 £ T £ 1253K and
5´10-11 £ P £ 2´10-2atm. S2
2.2.4 Effects of the sulfur pressure, temperature and gas
By using the point defect diffusion model based on the assumption that S2 (g) is the predominant gaseous species, we obtain the following relationship:
1 k= k ° P6 (2.19) r r S2
However, the point defect diffusion model used to predict the effects of sulfur partial pressure cannot be used to explain the behaviour of Fe1-dS very well. Instead, Young and Smeltzer [65] evaluated the integral in Equation (2.18) numerically, by using the data for Z [69] and the relationship between D and d and hence with P on the Fe Fe S2
T T assumption that DFe = D Fe (where D Fe is the tracer diffusion coefficient in the single crystal Fe S) for a number of P values. This yielded good agreement with the 1-d S2 measured result as shown in the Figure 2.8 [67]. This provides support for the applicability of Wagner’s theory to the sulfidation reaction.
The effect of temperature arises through the effect of P on the temperature S2 dependence of d and therefore on the effective valence appearing in Equation (2.18).
The activation energy for kr at constant sulfur pressure has been reported at values from
0 to 85kJ/mol according to the value of P . S2
In the study of sulfidation, P is often controlled by the following reaction: S2
H2S (g)=H2 (g)+1/2S2(g) (2.20)
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The presence of hydrogen has been assumed to have no effect on the sulfidation reaction, but this assumption has been questioned in interpreting the observation of porous sulfide [70, 71]. One suggestion is that dissolved hydrogen in iron sulfide becomes bound to iron vacancies and inhibits iron diffusion and enhances the sulfide ion mobility [70]. In another suggestion [71], hydrogen was considered as interstitial protons, and iron diffusion to be increased as the result of increased numbers of iron vacancies. However, in the view of the agreement between experimental kinetics and
Wagner’s theory, it can be concluded that hydrogen has no effect on the reaction, at least in the short term.
A duplex scale with a compact outer layer overlying a porous layer can form after extensive sulfidation of iron. A mechanism based on the fact that the metal surface recedes as it is sulfidized has been proposed [72, 73] and developed by Mrowec [74,
75]. It can be interpreted as follows: the metal sulfide is formed at the metal/gas interface by metal transport through the scale, and the gap between scale and metal left by the consumed metal cannot be filled by sulfide formation. To keep in contact with the metal, sulfide must therefore deform, presumably in plastic fashion. Separation from the metals occurs when the limit of scale plastic deformation capacity is reached.
As a result of the scale separation, the activity of sulfur at the scale under surface is raised and metal sulfidation continues. In fact, porous sulfides develop in the gap and form a bridge between the original scale and metal. Further growth of the outer scale is then supported by renewed outward metal transport. Growth of a porous underlayer is explained on the basis that the rate-controlling step is gas diffusion across the gap [76,
77], and irregularities in the surface of the newly forming sulfide are magnified under these conditions, leading to a porous structure.
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Figure 2.8 Sulfur partial pressure dependence of the parabolic rational rate
constant for the sulfidation of pure iron [67].
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In summary, it can be concluded that the early stage of the iron sulfidation is controlled by the lattice diffusion of iron through the growing scale. Mechanical failure at the scale/metal interface leads to the formation of a duplex scale when the extent of reaction is large [63].
2.2.5 Sulfidation of binary alloys
The general principles of binary alloy oxidation can also apply to sulfidation reactions. According to the Ellingham diagram, many alloying element sulfides are more stable than Fe1-dS. Such alloying additions form their sulfides at the place where the sulfur activity is lowest, e.g. either at the scale/alloy interface or beneath the alloy surface as an internal precipitate. Since internal sulfidation is harmful, it should be avoided. If the more stable sulfide is formed as an external scale, and it has a low diffusivity for iron (if iron is partially soluble in the other sulfide and tends to diffuse outwards though the more stable sub-layer) or the more stable sulfide can block the iron outwards diffusion (if the solubility of iron is negligible in the more stable sulfide layer), it will increase the sulfidation resistance of the alloy. The kinetics of internal sulfidation in the case that has no external scale and inward diffusion of sulfur is rate- controlling step are parabolic [50]. A variety of reaction rate laws are possible in the case of both internal sulfidation and external scale formed on the metal [52].
2.2.5.1 Sulfidation of Iron-Chromium
The effects of adding chromium to iron based alloys to promote their sulfidation resistances have been investigated [78]. It was reported that little change was produced in the sulfidation resistance rate with the addition of chromium up to 4wt%, but the sulfidation resistance was significantly increased at levels between 4 wt% and 12 wt%
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in the temperature range of 973-1073k at Ps2=0.066atm. It was found that chromium concentrated in the inner part of scale in the form of daubreelite, FeCr2S4 [79].
An investigation of the sulfidation properties of an Fe-26.6at%Cr alloy at temperatures of 973-1173K at P = 10-4-10-6atm was undertaken by Narita and S2
Smeltzer [80]. The sulfidation kinetics exhibited an early-stage transient interval, followed by a long-term stage of parabolic kinetics during which metal migration through the scales was the rate-controlling step. Different scales were formed, depending on the sulfur partial pressure: a triplex (CrFe)Sx/(CrFe)3S4/-(FeCr)Sx scale formed at P ³10-2Pa; a single phase (CrFe)S or a duplex (CrFe)S /(FeCr)S formed in S2 x x x the range of 10-2Pa³ P ³10-5 Pa and a single phase (CrFe)S formed when P was S2 x S2 lower than10-5 Pa.
In the study of Mrowec and Werber [81], it was found that, at a P of 1atm and in S2 the temperature range of 973-1273K, all alloys containing from 0.35 to 74 at% Cr sulfidised according to parabolic kinetics. Figure 2.9 [82] shows the variation of rate constant with alloy chromium content. In region I, the scale consisted of Fe1-dS layer and small amounts of spinel but no dissolved chromium. In region II, with increasing levels of chromium, the rate decreased because the displacement of iron with chromium in the mixed sulfide led to a low diffusivity of iron in this phase, and the growth of the outer layer was slowed down. The scale in this region was duplex, including an Fe1-dS layer containing some dissolved chromium overlying a layer of Fe(Fe2-xCrx)S4. In region III, a single layer of Cr2S3 containing some dissolved iron formed. Comparing the sulfidation rates in this region with the results for pure chromium, it was shown that the former were even slower than the latter. Marker measurements showed that scale growth was due solely to outward diffusion of both metals (Figure 2.10) [75]. By 35 Literature Review
invoking the Wagner-Hauffe doping theory [83, 84], the result was explained by
2+ assuming that Fe was incorporated into a metal deficit Cr2S3 structure. The vacancy concentration was thereby reduced, and if the diffusion in Cr2S3 proceeded via cation vacancies, then the slower rates are reasonable. But Cr2S3 is not a metal deficit sulfide
[59, 64], so it was concluded that chromium interstitials are the mobile defects.
Narita and Nishida [85, 86] found that the sulfide formed on high chromium alloys was Cr3S4, but not Cr2S3. The doping effect of iron on the growth of this sulfide is much easier to understand, since it is believed that the main lattice defects in Cr3S4 are cation vacancies. Apart from these differences, the results for a range of Fe-Cr alloys reacted at P =1atm and 973£T£1173K in the work of Narita and Nishida were in good S2 agreement with those of Mrowec and Wcber [81].
A series of Fe-Cr alloys was sulfidised in H /H S atmosphere ( P =2´10-11atm) at 2 2 S2
813K by Zelanko and Simkovich [87]. Parabolic kinetics were found and the dependence of kp on alloy chromium content was close to that found in S2 (g) at 1 atm.
However, a slight increase of kp was found for Fe-1wt%Cr alloy compared to pure iron.
This was explained on the basis of the following doping effect:
x × " (2.21) 2Cr + 3FeFe = 2CrFe +V Fe + 3Fe
Obviously, it leads to a higher concentration of vacancies and an increased the diffusion rate.
It is clear that the diffusion of metals through the growing scale is the rate- controlling step in the sulfidation of Fe-Cr alloys. The partially protective sub-layer of chromium sulfides is formed next to the alloy surface because of the greater stability of chromium sulfides than Fe1-dS. Compared with Cr2O3, chromium sulfides are more
36 Literature Review
conductive and faster growing than Cr2O3, and sulfur, unlike O2, fails to passivate the alloy surface due to the high lattice defect densities of the chromium sulfides.
Figure 2.9 Sulfidation rate constant for Fe-Cr alloys at P =1atm and T= 1173K S2
[82].
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Figure 2.10 Cross-section of scale formed on Fe-14Cr alloy reacted with P at 1173K S2
[75].
2.2.6 Sulfidation of commercial alloys
Since sulfidation of steels influences the equipment life severely, the choice of appropriate material is of economic significance. Suitable commercial steels must have reasonable sulfidation resistance and an acceptable cost.
2.2.6.1 Sulfidation of carbon steel
Nishida investigated the sulfidation of a series of steels with different carbon content (0.09, 0.56 and 0.95wt%) and also including 0.3wt% Si and 0.4wt% Mn at temperatures of 773K to 1173K in sulfur gas. The results were compared with those of
Armco iron. The kinetics of scale growth were identified as parabolic. With increasing carbon content, the rate was slightly decreased [88]. The reason was thought to be the barrier effect of graphite precipitates. Sulfidation kinetics of carbon steel containing
0.06wt% C have been reported as parabolic initially, followed by an approximately
38 Literature Review
linear rate [89]. The linear kinetics were due to formation of a porous outer scale consisting of Fe1-dS, leading to rapid inward diffusion of gas.
2.2.6.2 Sulfidation of low alloy steels
The sulfidation behaviour of low chromium and Cr-Mo steels, which are known to resist attack by organic sulfur compounds in crude oil, were reported to be unsatisfactory in H2/H2S atmosphere by several investigators [90, 91]. The sulfidation rates of steel containing less than 10 wt% chromium were found to be the same as that of plain carbon steel by Sorell and Hoy [90]. These results were in coincidence with the works of Sorell [92] and Backensto [91]. It was also shown that with increasing chromium content, the sulfidation rate was slightly decreased [89, 91]. An investigation of the variation in sulfidation rate with H2/H2S ratio and temperature for steel containing up to 5wt% Cr was undertaken by Backensto and Sjoberg [93]. The sulfidation mechanism of low chromium steels was proved to be different from that of pure iron, although the approximate same rate was found. The scales formed on these steels were found to be bulky, porous and brittle, and hence prone to spallation.
2.2.6.3 Sulfidation of ferritic stainless steels
The sulfidation of ferritic stainless steels has been investigated by many investigators [90, 92, 94]. Steels with chromium content about 12wt% have been studied in H2/H2S at temperature about 900K, and their sulfidation resistance was twice that of carbon steel in a similar environment [90-92]. Figure 2.11 shows that the sulfidation rate of those steels with higher chromium levels are slowed rapidly. In the study of Sorell and Hoyt [90], the sulfidation of the stainless steel Types 405 and 410 were found to exhibit irregular kinetic behaviour. This can be explained by using
Figure 2.11. These steels contain 11.5 to 13.5wt% Cr, which are in the rapidly
39 Literature Review
changing portion of the sulfidation rate curve in the Figure 2.11. Mrowec et al [94] showed that the relationship between sulfidation rate and steel chromium content depends on the temperature. In the above studies, it was also reported that irregular sulfidation kinetics were found for stainless steels containing 13 to 25wt% Cr at 873 to
1273K under P =1atm. However, parabolic kinetics were found when the temperature S2 was above 1273K. The scale consisted of an Fe1-dS layer overlying an intermediate layer of mixed sulfide overlying an inner layer of Fe1-dS, Cr2S3, FeCr2S4, MnS and SiS 2.
2.2.6.4 Sulfidation of Austenitic stainless steels
Bruns [95] investigated 200 series steel 16Cr6Mn4Ni and 17Cr8Mn4Ni at 866K in pure H2S, and found them following the same sulfidation rate as 16Cr steel. The 300 series Cr-Ni steels have been studied by many workers [90, 96-97]. Conflictive results were obtained for the sulfidation of types 304 and 309 steel. Bruns [95] found that the sulfidation resistance of 304 was same as for one chromium steel with the same chromium content. However Sorell and Hoyt [90] considered that the sulfidation of 304 was better than the corresponding ferritic chromium steel [90, 95]. The 304 steel was at least 10 times as resistant as plain carbon and low alloy Cr-Mo steel in sulfidation experiments [90]. Comparison between Types 309 and 304 has been undertaken by
Bruns [95] and Backensto et al [93]. The former described the sulfidation rate of the
309 steel as a little slower than that of the 304 steel, and the latter considered there was no difference between these two steels. They agreed that Type 310 has better sulfidation resistance than both types 304 and 309.
The effects of temperature and the role of chromium on the sulfidation kinetics of
310 stainless steel have been studied in 1.4´10-9< P <3.7´10-4 atm at temperatures S2 from 910K to 1285K [96-97]. The weight gain results showed that kinetics were 40 Literature Review
parabolic after an initial transition period. In all cases, the scale consisted of a duplex layer. X-ray diffraction analysis and EPMA were employed to characterise the scale.
These results showed that the outer layer contained phases of the Fe-Ni-S system, and the inner layer was a Fe-Cr-S layer, overlying an internal sulfidation zone of chromium sulfides and FeCr2S4 (Figure 2.12 [97]).
Figure 2.11 Sulfidation rates of chromium steels in H2S/H2S at 12 atm at 908K
[91].
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Figure 2.12 Scale formed on type 310 steel at P =1x10-7 atm and 1145K [97]. S2
The sulfidation of steel containing 23Cr-13Ni and 23Cr-18Ni over the temperature range from 973K to 1373K and P =1atm has been investigated [98]. In each case, a S2 duplex scale consisted of an outer layer of (FeNi)1-dS rich in iron overlying a layer containing Cr2S3, FeCr2S4, NiS and FeS. The short-term sulfidation kinetics were identified as parabolic. The higher nickel content was associated with a higher sulfidation rate. Scale growth was attributed to the outward diffusion of metal ions.
It is clearly that the sulfidation mechanisms of alloys have not been well established. The addition of chromium and aluminium has been shown to be efficient in promoting the sulfidation resistance of steel. However, it is still not possible to use the classical theory to explain their sulfidation behaviour since important detailed information on their solid-state chemistry is still lacking [63].
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2.3 HIGH TEMPERATURE CORROSION OF METALS AND ALLOYS IN
ATMOSPHERES CONTAINING HCl/Cl2
2.3.1 Introduction
It is well know that the corrosion rate of metals can be accelerated in atmospheres containing HCl/Cl2, but the detailed mechanisms are not as clear as for oxidation and sulfidation [99]. Daniel and Rapp [100], and Haanappel et al [101] reviewed the effects of chlorine on high temperature metal corrosion in oxidising environments. A mechanism has also been proposed, describing it as “active oxidation”.
2.3.2 Thermodynamic aspects
If metal M is exposed in an oxidising atmosphere containing HCl/Cl2, besides metal oxidation, the following reaction also occurs (assuming M is a divalent metal):
M+Cl2® MCl2 (2.22)
The value of P is given by: Cl 2
a DGo (2.23) P = MCl2 exp( MCl 2 ) Cl 2 am RT where P is the partial pressure of chlorine, a and a are the activities of MCl Cl 2 MCl 2 m 2 and M, respectively, and DGo is the standard free energy of formation of MCl . The MCl 2 2 thermodynamic data related to iron, chromium and nickel have been reported by Ihara et. al. [102-106]. Most of the metal chlorides have relatively high vapour pressures. In contrast to sulfur and carbon, chlorine is insoluble in the metal phase. Therefore chlorine is more likely to concentrate at the metal/scale interface. The formation of volatile metal chlorides is then possible when the information pressures are exceeded.
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2.3.3 High temperature corrosion of iron in atmospheres containing HCl/Cl2
Many investigators studied the high temperature corrosion of iron in gases containing HCl, or Cl2 together with O2 [105-111]. The corrosion products may be:
solid FeO, Fe2O3, Fe3O4, FeCl2, FeCl3 and FeCl 2(g), FeCl 3(g). In gas mixtures
containing HCl/Cl 2, Ihara et al [105] reported approximately parabolic rates at temperatures up to around 450°C, corresponding to the formation of solid chloride scale. Above 450°C, chloride sublimation occurred. At temperatures up to 600°C, sample weight decreased according to linear kinetics. This was explained again, as the result of chloride sublimation. No scale was found at higher temperatures [112]. The related reaction was described by Jacobson and Fruehan et al [106, 113-115].
Fe+2HCl(g) ® H2(g)+FeCl2(s)®FeCl 2(g)+H2 (g) (2.24)
It was considered that transport of gaseous chloride was the rate-controlling step.
The chlorination of iron will be changed if O2 is added to the atmosphere. At 550°C, the formation of FeCl 3 was increased, and therefore the corrosion rate was accelerated.
Ihara et al [105] suggested that the following two reactions were involved:
FeCl2+HCl+1/4O2®FeCl3+1/2H2O (2.25)
1/3Fe2O3+2HCl®2/3FeCl3+H2O (2.26)
It is known that the vapour pressure of FeCl 3 is very high, and a small amount of O 2 addition to HCl would accelerate the corrosion rate. However, at higher O 2 partial pressure, iron oxides were formed and the scale became more protective. It was still porous, and not a satisfactory scale [106].
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2.3.4 High temperature corrosion of chromium in atmospheres containing
HCl/Cl2
According to a few workers [103, 116], the corrosion products in gases containing chlorine and O 2 may be: a) solid state Cr2O3, CrCl 2 and CrCl 3 and, b) gaseous CrCl 2,
CrCl 3, CrCl 4, CrO 2Cl2 and CrO 3.
In pure HCl, only CrCl 2 was formed [103]. At temperatures below 600°C, approximately parabolic kinetics were found. At higher temperatures, the sublimation of CrCl 2 led to linear weight loss kinetics.
At 700°C, the corrosion of chromium in the gas mixture ( P <5´10-2atm) was Cl2 studied by Reinhold and Hauffe [116]. The sublimation of CrCl 3 was considered as the main reason for linear weight loss kinetics. CrCl 4 was found at a higher Cl2 partial pressure. The above authors proposed that the reaction followed the following steps:
· reactants diffused through the laminar boundary layer and then the scale
· chromium chlorides formed according to the different temperature and gas
composition
· solid state corrosion products vaporised:
CrCl2(s)®CrCl 2(g) (2.27)
· the volatile products diffused across the laminar boundary layer
In gas mixtures containing Cl2 and O2, the formation of Cr2O3 and chromium chlorides happened together in the initial period, and subsequently chlorides could be overlayed by oxide [103, 116]. Whether or not the resultant scale was satisfactory or not depended on the oxygen content of the gas. The chlorination of chromium was studied at 700°C in the gas mixture (P =2.6´10-2 atm) with different oxygen partial Cl 2
45 Literature Review
pressure (2.6´10-3 atm, 2.6´10-2atm and higher) by Reinhold and Hauffe [116]. A weight loss was noticed at P =2.6´10-3 atm since there was no protective scale formed. O2
At P =2.6´10-2atm, a weight gain was observed due to the formation of chromia. O2
Reinhold and Hauffe [116] explained the reason for the weight loss at higher oxygen partial pressure as being due to the following reaction:
1/2Cr2O3(s)+Cl2(g)+1/4O2(g)®CrO2Cl2(g) (2.28) The effect of oxygen partial pressure on the chlorination of chromium has also been investigated [103] at temperatures up to 500°C with oxygen partial pressure range from
0.2 to 1atm, where the scale was reasonable protective. However at high temperature, the corrosion products failed to protect chromium from serious corrosion (see Figure
2.13). It is considered that the severe corrosion was due to formation of CrCl3 and water vapour:
1/3Cr2O3(s)+2Cl(g)®2/3CrCl3(g)+H2O(g) (2.29)
2.3.5 The chlorination of alloys containing chroumium
The oxidation and chlorination of 2.25Cr1Mo steel at 500°C and 25Cr20Ni steel at
700°C were investigated by Reese et al [117] in an N2-H2-H2O-HCl gas mixture. It was calculated from gas composition that Fe3O4 on the low alloy steel and FeCr2O4/Cr2O3 on the high-alloy steel were the stable scales at the oxygen pressures in these experiments. Chlorides could also be formed due to the high chlorine gas content.
Their results showed that simultaneous growth of oxides and chlorides occurred in both cases for the samples after 1 and 10 min exposure. After 10 min, the oxides had overgrown the chlorides and the thermodynamically stable stage was reached. Metal chlorides were found beneath the oxide scale. However, the formation of chlorides during the initial period significantly affected the oxide morphology: no protective
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oxide scale was obtained. This was attributed to several reasons: the chlorides at the metal/scale interface cannot be dissolved by the base metal; gaseous metal chlorides resulted in scale cracking and the formation of porous oxide scales and chlorination of metal may be continued if the HCl has continuous access to the metal/scale interface.
50% O2
20% O2
75% O2
100% HCl
Figure 2.13 Diagram of change of weight losses of Cr2O3 exposed to reaction gas for 2h
[103].
Bramhoff et al [102, 118-119] studied a 10CrMo9.1 steel in an atmosphere containing HCl gas (87%H2/13%O2/HCl) at levels ranging from 0 to 3000ppm at
500°C, and Fe-20Cr alloy at 900°C in an atmosphere with P =1´10-20 atm and HCl O2 partial pressure varied from 0 to 2000ppm. In the case of the low chromium alloy, weight gains were observed to continually increase with increasing HCl levels up to approximately 1000ppm. However, a weight loss resulted upon exposure to higher
47 Literature Review
levels of HCl. This was attributed to the evaporation of the corrosion product. The scale formed on the low chromium alloy was described as porous and consisting of
Fe3O4 and some Fe2O3. At the metal/scale interface, FeCl2(s) was found. Therefore,
FeCl2 (g) was believed to be the gaseous corrosion product at metal/scale interface. At higher temperature and lower O2 partial pressure, the evaporation rate was considered to play the dominant role. In the case of Fe-20Cr, samples weight gain decreased with increasing level of HCl. It was proposed that both oxidation and chlorination happened in the initial period. If Cr2O3 formed on the alloy, the reaction between chromia and
HCl led to evaporation [102, 119].
Cr2O3(s)+4HCl(g)+H2(g)®2CrCl2(g)+3H2O(g) (2.30)
A stationary state was reached, resulting in a linear weight loss. Gesmundo et al [120] also reported the observation of this steady state.
According to Strafford [121], it could be concluded that increasing alloy chromium content can control the chlorination gradually. In his study of the chlorination of Fe-Cr alloys containing 2, 5, 9, 14 and 25%Cr at 1000°C in an atmosphere with P =10-16atm O2 and P =10-8atm. Corrosion occurred through the grain boundaries in the initial Cl 2 periods. The corrosion products included FeCl2 and Cr2O3.
The effect of chlorine on mixed oxidant corrosion of stainless steels at 540°C was studied by Bakker [122]. To produce the complex conditions in coal gasifiers, chloride rich deposits (75-80 wt% fly ash, 10-20 wt% carbon black, 3-10 wt% FeCl3 and 2-5 wt% NaCl) were used. Thus local HCl concentrations were much higher than those calculated from the chlorine content of the coal. Details of the gas compositions are listed in Table 2.1. Alloy compositions investigated are shown in Table 2.2. The results showed that internal oxidation of chromium caused a Cr2O3-rich but porous scale
48 Literature Review
which was not protective. The formation of this undesirable internal oxidation was attributed to the fast removal of iron and nickel from the alloys as chlorides. These chlorides will decompose in contact with reaction gases to form a very porous (Fe,
Ni)xSy outer layer. Oxidants like H2O and CO2 could penetrate this porous scale to reach the metal/scale interface resulting in internal oxidation of chromium and other oxide formers. Therefore, the corrosion losses increased dramatically with increasing chlorine levels. Higher levels of chromium or other oxide formers like silicon or aluminium in the alloy will protect alloys from attack by chlorine. In the case of such alloys, a higher HCl content is required to trigger the corrosion mentioned above. The work also showed that at room temperature the migration of chloride-rich liquids to the scale/metal interface frequently resulted in the formation of aqueous corrosion products such as iron chlorides that can cause scale spallation during re-exposure, and lead to significantly increased corrosion rates.
2.3.6 Chlorination of alloys containing aluminium
As has been discussed, the addition of aluminium to alloys can improve both oxidation and sulfidation resistance. In environments containing HCl/Cl2, the corrosion resistance of aluminium alloys has been proved better than that of alloys containing chromium [123-125]. Obviously, the better corrosion resistance is due to the formation of Al2O3. Elliot et al [126] suggested that the following reactions occurred. This was proved by the presence of chloride at the alloy surface.
Al+3/2O2®1/2Al2O3 (2.31)
2Al+3C l2®2AlCl 3 (2.32)
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Table 2.2 Alloy compositions [122].
Ni Cr V Si Mn Fe C Other
A 34.9 20.8 0.4 0.2 Bal <0.1 -
B 34.5 20.2 3.6 0.3 1.1 Bal <0.1 -
D 34.4 20.8 - 3.7 0.9 Bal <0.1 -
Alloys 800 33 21 - 1.0max 1.5max Bal 0.1 0.4Al
0.4Ti
310 20 25 - 1.0max 2.0max Bal 0.25max -
Oh et al [127] reported that when superalloys were exposed to Ar/20%O2/2%Cl2 at
900°C, those containing sufficient aluminium or titanium sustained the smallest attack.
The alumina layer was considered to act as a barrier between the alloy and the environment. Beneath the scale, no chloride was found. Small amounts of SO2 had no obvious effect on the corrosion of metal in this case [125].
Because of the better chlorination resistance of aluminium-containing alloys at high temperature, Fe-Al alloys were developed and studied for service in chlorine-containing environments. A detailed review of the chlorination behaviour of Fe-Al alloys can be found in section 2.4.6.3.3.
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2.3.7 Active oxidation mechanism
To explain the chlorination results, the mechanism of “active oxidation” has been proposed [128]: (Figure 2.14)
1. transport of reactants through a laminar gas diffusion boundary layer;
2. diffusion of reactants through the oxide scale;
3. formation of volatile metal chlorides;
4. diffusion of chlorides through the pores in the oxide scale;
5. formation of chlorine and metal oxide caused by the reaction between the
chloride and metal oxidation;
6. diffusion of the reformed Cl2 to the interface of metal and scale.
According to the literature reviewed, it is clear that the presence of hydrogen chloride or chlorine increased the corrosion rate of metals and alloys due to the formation of volatile metal chlorides at the metal/scale interface. The corrosion products are less compact than those formed in oxidizing environments. The chlorination mechanism is more complex and therefore less well understood than those of oxidation and sulfidation. Detailed information related to the chlorination behaviour of metals and alloys under the experimental conditions of this project are still lacking.
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(1) Cl2(g) MCl2(g) MCl2(g) 1
Cl2(g) +MO(s)®MCl2(g) +1/2O2(g)
5 Cl2(g) (2) 2, 6 Cl2(g) +MO(s)¬ MCl2(g)+1/2O2(g)
4 M(s) + Cl2(g)®MCl2(s)®MCl2(g) 3 (3)
(1) Diffusion boundary layer, (2) oxide scale, (3) substrate.
Figure 2.14 Schematic overview of several reactions considering the induced
corrosion attack by the presenc e of HCl or Cl 2 [129].
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2.4 PACK CEMENTATION ALUMINIZED COATING
2.4.1 Introduction
The potential of iron-aluminide alloys as oxidation-resistant materials has been investigated for more than 60 years [129]. The excellent oxidation and corrosion resistance, low density and low material cost of iron aluminide alloys have already been widely recognized. However, it is also known that iron-aluminide alloys are limited for many structural applications by their poor mechanical strength [130, 131]. Therefore, the idea of applying an aluminized coating on materials which have adequate strength but inferior high temperature oxidation resistance has received considerable attention.
The pack cementation aluminized coating is one kind of high temperature coating considered. High temperature coatings have been studied by many investigators over the past 50 years [132-135]. They can be divided into the following three groups [136-
137]:
a) Diffusion coatings;
b) Overlay coatings;
c) Thermal barrier coatings.
Diffusion coatings have been applied by using chemical vapour deposition (CVD), pack cementation, and slurry cementation. Comparing with the others, the pack cementation coating has become widely used at present because of its low equipment cost, simplicity and flexibility in dealing with complex shaped components [138].
2.4.2 The pack cementation process
Pack cementation is one of the chemical vapour deposition methods. This process was first used in the deposition of aluminium on copper and iron in 1911 and 1914
54 Literature Review
[139]. Since then, this procedure has been widely investigated. Today, our understanding of the principles and kinetics are well developed and the procedure has been applied to a great range of substrates.
2.4.3 Procedure
The pack cementation procedure can be described as following: the uncoated substrate is packed in the cement which consists of intimately mixed powders of either pure or alloyed depositing elements (master alloy), a halide salt (activator) and an inert filler (mostly Al2O3) which is used to prevent the pack from sintering in a semi-sealed retort. The master alloy could be Al, Cr, Si, Cr-Al, Ni-Al etc. Typically, NH4X, NaX,
AlX3, where X refers to F, Cl, Br, I, are used as activators. The packed and semi-sealed retort is heated under flowing inert (e.g. Ar) or reducing (H2) gas (or a mixture of both gases) to the coating temperature, which is enough for the chemical reactions and the diffusion to proceed. A schematic diagram of the typical pack cementation arrangement is shown in Figure 2.15 [140]. Coating formation may be considered in the following steps [141]:
1) The formation of gaseous metallic halides (MXn) by reaction between
decomposed or vaporized activator and the master alloy. For example, in the
case where NaX is the activator salt, the following reactions may occurs:
2NaX(l)®2Na(g)+X2(g) (2.33)
M(s)+n/2X2(g)®MXn(g) (2.34)
2) Diffusion of the gaseous metallic halide species formed in the pack atmosphere
to the surface of the substrate driven by the activity or chemical potential
gradient.
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3) Deposition of the coating element takes place to decrease its activity. This step
includes: adsorption, dissociation, surface diffusion and a number of other
steps. The possible mechanisms of dissociation of metal halides at the substrate
are [142]:
a) Disproportionation reactions
2MX(g)®M+MX2(g) (2.35a)
3MX2(g)®M+2MX3(g) (2.35b)
b) Formation of salt (as melt) at substrate surface
MX(g)+Na(g)®M+NaX(l) (2.36)
c) Reaction of species within gaseous environment
MX(g)+1/2H2(g)=M+HX(g) (2.37)
4) Diffusion of deposited element(s) into substrate leading to coating growth. The
coating morphology and phase constitution depend on the thermodynamic and
kinetic factors such as the activity of the coating elements in the gaseous phase,
stability of the formed phases and the diffusivity of the elements in these phase
2.4.4 Mechanism of coating formation
Coating formation is governed by the alloy composition, in fact, the activity of the aluminium in the master alloy [143-145]. The activity of Al in the pack determines what kind of aluminide phase is stable at the interface between the coating and the pack.
Table 2.3 lists the Al activities in Fe-Al system at 900°C studied by J. Eldridge et al
[145]. A phase diagram of Al - Fe [146] is shown in Figure 2.16. By controlling the Al activities in the master alloy, coatings with different desired Al compositions can be produced.
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. ] 0 Schematic diagram of a pack cementation retort [14 Figure 2.15:
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Table 2.3 Aluminium activities in Al-Fe system at 900°C [145].
Al (at%) aAl gAl
5 3.8 x 10-4 7.6 x 10-3
10 8.1 x 10-4 8.1 x 10-3
15 1.5 x 10-3 9.8 x 10-3
20 2.45 x 10-3 1.2 x 10–2
25 3.7 x 10-3 1.5 x 10–2
30 5.5 x 10-3 1.8 x 10–2
35 8.3 x 10-3 2.4 x 10–2
40 1.35 x 10–2 3.4 x 10–2
45 2.6 x 10–2 5.75 x 10–2
50 5.5 x 10–2 1.1 x 10-1
52 9.8 x 10–2 1.9 x 10-1
60 9.8 x 10–2 1.6 x 10-1
67 1.3 x 10-1 1.9 x 10-1
70 1.3 x 10-1 1.8 x 10-1
75 4.0 x 10-1 5.2 x 10-1
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[ 146 ] . Figure 2.16 Fe - Al binary alloy phase diagram
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2.4.5 Process variables in pack cementation
There are a number of important variables that can affect the formation of the pack aluminized coating through their influence on thermodynamic and kinetic factors.
Some of them are [147]:
· Chemistry of the substrate and the pack mixture
· Quantity of activator in the pack
· Purity of the powder
· Homogeneity and permeability of the pack mixture
· Variables of the procedure (pressure in the retort, temperature and deposition
time)
2.4.6 Pack aluminizing on Ni-based and Fe-based alloys
2.4.6.1 Pack aluminizing on Ni-based alloys
Although pack aluminizing of Ni-base alloys is not the major interest in this project, the extensive work carried out in this area has yielded an understanding of the process basis.
Seigle et al [132] studied the effect of varying Al activity in the pack on the surface composition of aluminized coatings on pure Ni. Atmospheres of flowing pure H2, Ni-
Al master alloys (50-70 at% Al), and AlF 3 or NH4HF activators were used in the experiments. The results showed that solid-state diffusion was the rate-controlling step when the Al activity was low. The surface composition of a coating produced in this way was below 50 at% Al. In the case of high Al activity, the rate-controlling step was a combination of solid and gaseous diffusion. A steady state was reached, although the equilibrium did not exist at the interface. The Al activity value of the steady state was found to be below that of the master alloy. This supported Levine and Cave’s model
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[148] for gaseous diffusion, in which a constant Al activity at the coating surface was assumed in the thermodynamic calculations.
In Seigle’s subsequent works [133-135], the effects of different activators: AlF 3 or
NaF, NaCl or NaI on the aluminizing potential of the pack were investigated. Either pure Al or Ni-Al alloy was used as the master alloy to be deposited on a pure Ni substrate in flowing pure H2 over a range of temperatures. The largest weight gains were found in the case of fluorides as the activators. The ranking of the efficiency of the investigated activators was AlF3>NaF>NaCl>NaI. Parabolic kinetics were found when fluorides or chlorides were the activators. In the case of NaI as the activator, low and erratic weight gains were obtained. Some of the results are shown in Figure 2.17
[134]. These results are coincident with those of Levine and Caves [148]. It was assumed that the chemical reactions occurred so quickly that local equilibrium could be established at the pack-coating interface. The halide vapour pressures at the substrate surface were therefore assumed to be at a constant aAl, determined by mass balance and thermodynamic equilibrium conditions at the surface. This established a boundary condition for the mass transfer process.
When Al is transferred from the pack to the surface of the substrate, the width of the depleted zone and the diffusion distance for the halide vapor increase. The parabolic relationship is used to express the amount of Al (Wg) transferred by diffusion
(in g ×cm-2):
2reM N d (2.38) W 2 = Al t = k t g l A g where r is the pack Al concentration(g×cm-3), e is the correction factor for pack porosity and l is the pore length, M is the Al atomic weight, d is the diffusion distance (cm), A is
2 the total area (cm ) and NAl is the overall rate of diffusion of Al through the gas phase
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4 wt% AlF3
4 wt% NaF
4 wt% NaCl
4 wt% NaI
Figure. 2.17 Variation of Ni-specimen surface composition and weight gain with time
in 4 wt% aluminium packs at 1000°C [134].
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(mol sec-1), which can be calculated by the following equation:
N Al d 1 ' (2.39) = å Di - PP ii )( A RT i
’ here Di is the diffusivity. Pi and Pi are the partial pressures of the ith Al-bearing species in the pack and at the surface of the coating, R is the gas constant and T is the absolute temperature. Combining the above two equations:
2reM ' (2.40) kg = å Di - PP ii )( lRT i
The calculation details depend on the type of the activator. In the case of Na (F, Cl,
Br, I) and AlF3 (1-10wt%), which are considered as condensed phases in the pack at the heat treatment temperature with flowing H2 or inert gas (argon), the following chemical reactions may occur:
NaX(l)+Al(l)=AlX(g)+Na(g) (2.41)
2NaX(l)+Al(l)=AlX2(g)+2Na(g) (2.42)
3NaX(l)+Al(l)=AlX3(g)+3Na(g) (2.43)
NaX(l)+1/2H2(g)=HX(g)+Na(g) (2.44)
NaX(l)=NaX(g) (2.45)
2AlX3(g)=Al2X6(g) (2.46) subject to the conditions,
PPPPPP= 3+ 6+ 2 + + (2.47) Na ()g AlX3 ()g Al2X 6 ()g AlX 2 ()g AlX HX ()g and,
SPi=1 (2.48)
’ The calculations of the Pi at the coating surface are similar, but the kinetics should be taken into account. Since the Na (g) and halogen must be transported in equivalent amounts, equation (2.47) leads to:
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' ' (2.49) D NaNa PP Na )( =- 3DAlX ( AlX PP 'AlX ) +- 2DAlX ( AlX PP AlX ) +- 3 3 3 2 2 2 D PP ' +- D PP ' )()( +- 6D ( - PP ' ) AlX AlX AlX HXHX HX Al X Al X 6262 Al X 62 )
For AlF 3 or NH4F activators, without the effect of Al activity, the following reactions can be used to calculate the partial pressure.
AlF 3(g)+2Al=3AlF(g) (2.50)
2AlF 3(g)+Al=3AlF 2(g) (2.51)
AlF 3(g)+3/2H2(g)=3HF(g)+Al (2.52) subject to the condition,
(2.53) å Pi = 1
considering the H2 balance, equation (2.53) can be modified to
2D (PP-' )() +D PP -' = 0 (2.54) HHH2 2 2 HF HF HF
For non-stable, or volatile activators such as NH4Cl (Br, I), dissociation occurs:
NH4Cl(g)=NH3(g)+HCl(g) (2.55)
2NH3(g)=N2(g)+3H2(g) (2.56)
By assuming the total pressure is 1atm, and the ratio of H to Cl is constant at 4:1, the partial pressures were worked out using the following reactions:
Al+HCl(g)=AlCl(g)+1/2H2(g) (2.57)
Al+2HCl(g)=AlCl 2 (g)+H2(g) (2.58)
Al+3HCl(g)=AlCl 3(g)+3/2H2(g) (2.59) subject to the conditions,
2 4( 32 +++=+ PPPPPP ) (2.60) H 2 HCl AlCl ALCl2 AlCl3 HCl
SPi=1 (2.61)
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Part of the results calculated by Seigle et al. are shown in Table 2.4 [135]. They are in good agreement with those of Levine and Caves [148]. In the calculation of the latter,
AlN(s) was taken into account when NH4X was the activator.
It should be pointed out that [141, 148] the efficiency of volatile activators ( e.g.
NH4Cl ), depends more on the initial amount of activator used than the temperature, whereas for the stable activators, the efficiency of the system largely depends on the temperature. The reason for the NH4Cl behaviour was the disappearance of the condensed phase, which maintains the uniform Al deposition capability, during the coating process. The other result was the finding that the fluoride activator was more efficient than the NaCl or NaI activator because of the higher aluminium halide partial pressures.
The kinetic models were developed by Levine and Caves [148]. A number of investigations have been undertaken by Seigle and co-workers [132-135, 149]. Most of the work was focused on the aluminizing of Ni and Ni-base substrates. Levine and
Caves [148], using pure Al master alloy and either stable or volatile activators (NaX or
NH4X) formed aluminide coatings on the superalloy IN100. The weight gain results followed parabolic kinetics in most cases. A depletion zone was found in the pack, in agreement with the observation of Brill-Edwards and Epner [150]. In calculating the net flux of Al toward the specimen surface, two distinct models for Al transport from the bulk pack to the surface of the coating were developed. The first, shown in Figure
2.18a [148], was used to explain the system in which volatile-type activators were used.
Since there was no condensed phase in the packs, the higher order Al halides diffuse back into the pack where they react with the aluminium source material. Consequently, the activator was circulating between the bulk pack and the substrate surface without being consumed. Figure 2.18b shows the other mechanism when stable activators are
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Table 2.4 Equilibrium partial pressure of gases in packs activated by AlF 3, NaF, NaCl
and NH4Cl at 1000°C [135].
Gaseous Pi (atm) aAl Species AlF 3 NaF NaCl NH4Cl AlF 1.0 5.18 x 10-2 2.69 x 10-2 - - -2 -4 AlF 2 " 1.09 x 10 6.31 x 10 - - -3 -3 AlF 3 " 9.51 x 10 1.32 x 10 - - HF " - 1.16 x 10-5 - - AlCl " - - 1.17 x 10-3 4.89 x 10-2 -4 -l AlCl2 " - - 1.08 x 10 1.37 x 10 -7 -2 AlC13 " - - 4.89 x 10 2.37 x 10 HCl " - - 1.11 x 10-5 5.83 x 10-4
H2 " - - - 0.7896 Na " - 3.21x10-2 1.41x10-3 - Ui (atm) AlF 10-2 2.40 x 10-3 1.85 x 10-3 - - -3 -4 AlF 2 " 2.35 x 10 2.98 x 10 - - -3 AlF 3 " - 4.28 x 10 - - HF " - - - - AlCl " - - 6.15 x 10-5 3.57 x 10-3 -5 -2 AlCl2 " - - 2.98 x 10 7.30 x 10 -7 -2 AlCl3 " - - 7.13 x 10 9.24 x 10 HCl " - - - 4.26 x 10-3
H2 " - - - 0.7685 Na " - 4.67 x 10-3 2.68 x 10-4
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(a)
(b)
Figure 2.18 (a) Activator “circulation” and
(b) “condensation” models for Al transport [148].
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Figure 2.19 Mixed mechanism of Al transport in the presence of both activator-
only, and activator and Al depleted zones [149].
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employed. In the case of AlF 3 as the activator, AlF3 (g) was transferred to the coating surface and then condensed as a solid phase. Therefore, the diffusion of Al occurs due to the condensation of the activator at the surface of the coating.
By using the principles of the gas phase model and solid-state considerations for the
Ni-Al coating system, Seigle et al developed a model of pack aluminizing process(see
Figure 2.19) [132-135, 149]. It is shown in Figure 2.19 [149], that aluminium can be transported to the substrate surface by a combination of the two mechanisms of Levine and Caves [148]: activator condensation and activator circulation. The first describes the transport of Al from the pack outside the depletion zones to the substrate surface where condensation of the activator is dominant. The second is used for Al transport from the activator-only depleted zone. This model was studied by comparing predicted and experimental results and found to be satisfactory.
2.4.6.2 Pack aluminizing of iron and commercial steels
Kung and Rapp [151] investigated the kinetics of iron aluminization using the pack cementation technique. In their study, the coating mechanism was investigated by using different halide activators and by tumbling the pack. The reason for tumbling the pack was to eliminate any chemical depletion zone. A porous alumina sheath was used to wrap the iron substrate to prevent the incorporation of powder particles into the surface.
Fe-Al alloy (40 at% Fe+60 at% Al) powder was used as the master alloy. NaCl, AlF 3 and AlCl3 were chosen as the activators. The coating process temperature was 900°C, and coating time varied from 5h to 40h. The method for analysing the thermochemistry was somewhat different from those used by Seigle et al [132-135]. By using the program SOLGASMIX [152], the bulk pack equilibrium gas phase composition was calculated. The results were in agreement with previous workers [132-135]. The effect
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of H2 was negligible: no significant differences were observed between the inert and reducing gas. They also found that AlF 3 was the superior activator in terms of aluminizing potential for iron substrates. In the case of NaCl as the activator, the experimental data were in good agreement with those predicted using SOLGASMIX
[152]. For the AlF 3 activated pack, the pack powder plugged the alumina wrapping sheath, which led to less aluminium being deposited on the substrate surface. The extremely volatile AlCl3 escaped from the pack, resulting in the loss of the aluminium source.
Pure Al powder (10-35wt%) was used as the master alloy in the work of Heckel et al.[153] Pure Fe and NH4Cl were used as substrate and activator, respectively. The
formation of a coating co nsisting of FeAl2, Fe2Al5 and FeAl was verified by X- ray diffraction (XRD). The thickness of each phase layer was measured after the reaction: all the data showed that the growth kinetics of the iron aluminides were parabolic.
Theoretically, the mechanical properties of these iron aluminides are quite similar to those of the bulk materials, which mean they are too brittle for many applications.
2 The microhardness values of FeAl3 and Fe2Al5 are around 1000kg/mm , but for FeAl, it is only about 600kg/mm2. For this reason, to lower the activity of aluminium in the pack, mixed iron aluminium powders sometimes were chosen as the master alloys.
Sivakumar and Rao [154] applied both pure Al powder and Fe-Al alloy (50wt%
Al+50wt%Fe) in the pack cementation process on mild steel, EN-3 (0.17% C).
NH4F. HF was used as the activator, and the processing temperature ranged from 750°-
900°C, with time varying from 2 to 20h. In the case of pure Al master alloy, FeAl3 and
Fe2Al5 were formed on the steel. However, the observed coating formation kinetics differed from those reported by Heckel et.al [153]. The weight gain kinetics obtained by Sivakumar and Rao [154] were found to be not parabolic, and this was attributed by
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them to the formation of the above two layers. Unfortunately, they did not give any further explanation of their hypothesis. For the coatings formed using Fe-Al master alloy, FeAl was identified using XRD. Alumina needles were found on the coating surface, which was attributed to the reaction between aluminium fluorides and H2O vapor in the pack. Coating kinetics were, in all cases, parabolic when an Fe-Al masteralloy was used.
The oxidation resistance of aluminium-rich coatings was also studied [154], and compared with 304-grade stainless steel at 900°C. It was found that, after 24h exposure in air, coatings formed in a short time (2h) had better oxidation resistance than did coatings grown for 10 and 20hr. This was explained in terms of the better coating structure: thinner coatings showed no porosity on the surface, but with the thicker coatings, porosity was more significant. However, thicker coatings showed better oxidation resistance after long time exposure than thinner coatings. Sivakumar and Rao
[154] also reported that for the coating formed at 850°C after 4h, all aluminium sources in the coating were consumed in a 72hr oxidation experiment. At this point the substrate started to be oxidized, and a very large weight gain resulted. The oxidation products were identified as a-Al2O3 and Fe2O3. The comparison material SS304 showed less oxidation resistance than did coatings grown for 10hr. Metallographic cross-sections of those coatings revealed (from marker behaviour) that they had grown by outwards diffusion, no matter what kind of master alloy was used.
The pack aluminizing process was also applied to a heat resistant stainless steel
MO-RE1 to improve its relatively poor high temperature corrosion resistance [155].
Kim et al used both pure Al powder (high activity process) and alloy powders of Fe-Al,
Ni-Al and Cr-Al as masteralloy. In most cases, NH4Cl powders were chosen as activator, other halides were also used occasionally. The microstructures of the coating
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layer were carefully controlled in order to obtain either inter-diffusion zone (mixture of a-ferrite and b-aluminide) or internal-diffusion barrier (mixture of s- and b-aluminide).
Pack aluminizing processes were carried out at temperatures from 900°C to 1050°C.
According to the results of SEM, EPMA and XRD, it was found that the structures formed in both internal-diffusion barrier and inter-diffusion zone formed on MO-RE1 were combinations of b-NiAl and a matrix of s phase (an Fe-Cr intermetallic compound stable at temperature up to 815°C [156]). It should be pointed out that s phase formed in the interdiffusion zone was transformed from a ferrite during cooling [155]. High temperature cyclic corrosion tests (45mins at testing temperature and 15mins at room temperature) were undertaken in the corrosive gas mixture (75%N2, 12%H2O, 9.5%CO2,
3.5%O2 and 200ppm SO2) at 1100°C. Results showed that coatings with the internal- diffusion barrier had much better high temperature resistance than those with the interdiffusion zone, in cyclic corrosion tests. Furthermore, their high temperature resistance increased with the increasing of the coating thickness [155].
2.4.6.3 High temperature corrosion resistance of iron aluminide
2.4.6.3.1 High Temperature Oxidation of Iron Aluminides
The excellent oxidation resistance of iron aluminides depends on the formation of a thermodynamically stable Al2O3 layer when exposed to oxidizing environment.
Different modifications of Al2O3 are formed [157]. The g, d, q and a phases contain different amounts of H2O and, of course, have different structures. Only the a phase is thermodynamically stable, the others being metastable. Therefore, it has been assumed that the same sequences of phase development are followed in high temperature oxidation of aluminium containing alloys. However, it has been recognized [158] that
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the formation of different modifications is mainly controlled by temperature and by epitaxial relationships. At low temperatures, metastable, fast growing q-Al2O3 is formed. High rates of q-Al2O3 formation were observed in the low temperature range up to 900°C [159]. The morphologies of this alumina oxidation product were described as needles and platelets. Between 900°C and 1000°C, the kinetics of iron aluminide oxidation were described [159] as a combination of fast oxidation rate in the initial stage, followed by a slow growth stage of scaling. At about 1000°C, the stable a-phase started to be the dominant oxidation product. A transition from q -Al2O3 to a-Al2O3 exists about this temperature [159]. At 1100°C, only a-Al2O3 could be found in the scale. Typically, a-Al2O3 is described as having a ridged structure [159].
Kolarik et.al [160] studied the oxidation products and their modifications on Fe-Al alloys containing 5-30 at% Al at 900°C in air with in situ high temperature X-ray diffraction method. Their results showed that at least 10at% Al was needed to form the protective a-Al2O3 scale on Fe-Al alloys at 900°C. To eliminate the formation of Fe2O3 completely, more than 20 at% Al was necessary. No q-Al2O3 or other alumina modifications were detected by in situ diffraction. However, deviations in the kinetics of a-Al2O3 were found: on Fe-20Al, the linear growth of the a-Al2O3 peaks differs from their growth on Fe-30Al and Fe-10Al. Further investigations on these phenomena observed in the high temperature oxidation of Fe-20Al are suggested. On the other hand, the method of using diffracted beam intensity changes of corrosion products to
deduce their formation kinetics is limited by the X - ray penetration distance into the oxide. Beyond this distance, the intensity of the same corrosion product can not be used as a measure of its accumulation.
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Beside the effect of temperature on the formation of different modifications of alumina, the aluminium level in the iron aluminide alloys is another key parameter in their oxidation resistance. Boggs [161] studied the oxidation behaviour of a series of iron aluminide alloys (to 8% Al) containing 0.1% C over the temperature range 450°C to 900°C. For his low temperature oxidation experiments, it was shown that the alloy with low aluminium content (0.003-0.095wt%) at 500°C had a faster oxidation rate than that of pure iron. Similar results also were obtained by Bateman and Rolls [162] although a different experimental environment was used. Bateman and Rolls [162] used a simulated burnt town gas mixture with or without 5vol% free O2 addition, while
Boggs used O2 at a pressure of 700 torr. The fast oxidation rate was explained as a doping effect. Trivalent Al 3+ replaced the Fe2+ in the scale, and increased the cation vacancy concentration and thereby the rate of diffusion of Fe2+ through the magnetite scale.
Boggs showed that the scale formed on alloys with (0.5wt% - 0.7wt%Al) at 500°C was made of iron oxides only, and the oxidation rate dropped to the same value as pure iron. When the aluminium level was increased to more than 1wt%, the oxidation rate was further decreased due to the formation of FeAl2O4 which acted as a diffusion barrier. For alloy with 4wt%-5wt% aluminium, the maximum oxidation rates were reached at about 570°C and decreased to minimum oxidation rates at 850°C. g-Al2O3 gradually became the dominant oxidation product with increasing temperature above
570°C [163-167], and started to decrease at temperatures above 900°C. At the same time, the amount of a-Al2O3 increased. Boggs and Hagel [163] believed the transition from g-Al2O3 to a-Al2O3 resulted from the recrystallisation of g-Al2O3.
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Iron-Silicon-Aluminium alloys were also investigated by Boggs [168]. The aim of using silicon as the second oxygen getter in iron-aluminium alloys was to suppress the internal oxidation of aluminium which would prevent or delay the rehealing of protective aluminium oxide layer [168]. Alloys containing 2-4% Al and about 2% Si were initially tested by means of microbalance in wet or dry oxygen and in water vapour at 800°C, and at 900°C in dry oxygen. The results showed a better oxidation resistance than those of Fe-5%Al alloys. To determine the optimum concentrations of
Al and Si, further experiments were undertaken. Alloys were exposed to flowing air at ambient humidity and to flowing air saturated with water at 32°C for 50 hours at
1093°C. An alloy with 6%Al and 1%Si was proven to have the best oxidation resistance among the alloys tested [168].
In the review of Tomaszewics and Wallwork [169], it was suggested that at least 8 wt% Al is required for the formation of protective alumina at 800°C [169-170]. The addition of Cr will reduce this level, while Ni will increase it. The Fe3Al and FeAl phases obviously exceed this critical aluminium level, and alumina is predicted to steadily form on these alloys above approximately 500°C when oxidized [130].
However, due to the thinness of the alumina and its susceptibility to spallation, it is very difficult to analyse this situation [171-172]. For example, large amount of scales formed on Fe-40at% alloy at 900-1100°C spalled after cooling, but it was also found that the spallation resistance tended to increase in cyclic oxidation experiments [173].
This phenomenon [173] was attributed to the spallation of small scale pieces instead of large ones. Consequently, less voids and bond-weakening segregants at the interface were formed. Parabolic rate constants for Fe-28at%Al alloy oxidized in dry air at
800°C were measured in the order of 10-5 mg2 cm-4 h-1 [171]. Small amounts of additional alloy elements like zirconium, hafnium and yttrium were found to have
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beneficial effects on the alumina scale formed on Fe-(8-19%) Al alloys [174-177]. For those alloys which have higher Al concentration, e.g. Fe3Al, further detailed investigations are needed. The oxidation behaviour of iron alloys containing 16-40% Al in air at 800-900°C were studied by Tortorelli et. al [178]. External alumina scales were formed to protect the substrate underneath. The measured rate constants were shown to be closer to those of a-Al2O3 growth than for other forms of alumina.
2.4.6.3.2 High Temperature Sulfidation and Oxidation/Sulfidation of iron aluminides
Comparing the sulfidation behaviour of iron aluminides to those of other iron based alloys (e.g. Fe-Cr alloys), iron aluminides have better sulfidation resistance [179-183].
This is because aluminium sulfide, Al2S3, is more stable than other metal sulfides thermodynamically, and has relative lower growth rate than that of iron or chromium.
Its large molecular volume results in a large Pilling-Bedworth ratio, so the scale is thought to be more protective) [184].
Sulfidation behaviour of iron aluminides containing more than 18at% Al were studied in H2S/H2 [180, 181, 183, 185]. Slow parabolic kinetics rate were measured
2 -4 -1 (£1mg cm h ) at 800°C with S2 partial pressure less than 1 Pa. Relatively good sulfidation resistance of those iron aluminium alloys was observed up to ~750°C. At higher temperature (800-1000°C) and with S2 partial pressure more than 133Pa, iron sulfide started to grow and a fast sulfidation rate was observed. Depending on the activities of S2 and Al, the surface morphology of these iron aluminium alloys in
H2S/H2 gas mixture were described as multilayer structures, which consisted of an FeS layer at the scale surface, an intermediate layer (Fe,Al)3S4, and Al2S3 at the alloy surface. Internal Al2S3, sometimes together with FeAl2S4, precipitates were also formed
-3 in alloy containing 18-20% Al at 750 and 900°C at S2 partial pressure of 10 Pa [183].
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For an alloy containing 28%Al, only external Al2S3 was formed. Through the marker investigations of Strafford et al [180] and Patnaik et al [183], it was found that the formation of external FeS scale was controlled by the outward diffusion of Fe and, the inner layer growth resulted from the inward diffusion of sulfur. Internal sulfide precipitates also caused the mechanical fracturing of the alloy surface.
Aiming at the requirement of coal conversion and combustion materials, Fe3Al and
-5 -20 FeAl were studied in the simulated H2S-H2-H2O (Ps2= 10 atm, Po2=10 atm) gas mixture of at 700°C and 800°C [171]. To identify the effects of Al concentration, a series of Fe-Al alloy with 25-40at%Al, 2-10at%Cr, and 1-2at% Nb/Mo were produced by using arc melting. Thermogravimetric analysis (TGA) was employed to study their sulfidation kinetics. For Fe-28at%Al alloy exposed to the gas mixture for over 290 hours at 800°C, fine surface scratches from specimen preparation were still clear, and
XRD results showed only g-Al2O3 was formed. As it is well known that small additions of Cr can improve the room temperature ductility of Fe3Al based alloys, Fe-28at%Al alloys with 2, 4 and 10at% Cr were also tested for comparison purposes. For alloys with higher Cr levels, the weight gain was over 20 times that of binary alloys. No significant difference in weight gain was observed between iron aluminium with and without 2at%Cr. Although iron sulfide was stable at the activities of sulfur and oxygen used in the experiment, no iron sulfide was found, even if scale spallation occurred during testing. This was attributed to the faster nucleation and growth of Al2O3 compared to that of iron sulfide. High Al concentration was required to retain the good sulfidation resistance of iron aluminide alloys.
Tortorelli et al [184] applied an iron aluminium weld overlay (Gas Tungsten Arc) coating on the commercial steel 2.25Cr 1Mo and stainless steel 304 grade to study their
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sulfidation/oxidation behaviour in the gas mixture (5.4%H2S-79.4%H2-1.6%H2O-
13.6%Ar, in volume) at 800°C. The partial pressures of oxygen and sulfur were 10-22 atm. and 10-6atm., respectively. A 2-3 mm thickness coating was produced on the steel surface. Only weld materials themselves were investigated in this study. The final Al content in those coatings varied from 11.1-36.7at%. Thermogravimetric studies showed these overlay coatings fell into two groups, depending on their Al level. The higher the
Al level, the more protective was the coating. Scales formed on high Al coatings were thin and uniform. Only small amounts of powdery scale spalled from those samples. In contrast, significant weight loss was observed on low Al concentration samples, and the black corrosion products formed were relatively thick. SEM and energy dispersive x- ray analyses results showed those corrosion products consisted of iron and aluminium sulfides. The corrosion resistance of weld iron aluminide coatings with 35at% Al was equivalent to that of similar wrought iron aluminides. Weld overlay coating appeared to be a potential method to produce effective corrosion-resistance, but more investigations on the minimum Al level necessary to resist the sulfidation and oxidation were still needed. Since the above study was focused on the deposited materials themselves, the effect of the adhesion and interaction between coating and substrate were still not clarified.
The other iron aluminium (aluminium level ³ 20-25%) overlay technique GMA
(Gas metal arc) [185] was also applied on the same substrates (304 and 2.25Cr1Mo)
[184]. Corrosion tests were undertaken in the same gas mixture used in the previous study [184]. The weight gains of the overlay materials only were measured in thermogravimetric tests. All the weld overlays showed good corrosion resistance, including a GMA produced deposit with only 21at% Al. Two more GMA product deposits on 2.25Cr1Mo were tested in cyclic condition. Both samples were held 78 Literature Review
isothermally at 800°C for 72-74 hours and cooled down. One sample was reheated to
800°C for several more hours, and the other one was cycled to below 100°C and back to the exposure temperature twice. Significant increases in the weight gain rate were observed.
With the same gas mixture [184-185] at 800°C, DeVan and Tortorelli [186] tested three groups of Al-containing alloys which had Al contents ranging from 16-40at%: binary iron aluminium alloy, ternary Fe-Al-Cr (Cr contents up to 5%) and multicomponent Fe-Al-Cr alloys containing 0.05-0.1%Zr, 1%Mo, 0.5%Nb and trace amounts of carbon and boron. Some specimens were pre-oxidized in dry air for 1 week to help to identify the reaction mechanisms. Thermal cyclic sulfidation experiments were also carried out on selected samples. Two binary iron aluminium alloys with 28% and 35% Al content, respectively, were tested in the mixed gases initially. No significantly different weight gain was found on these two alloys. The weight gain of
Fe3Al in the gas mixture after 1 week exposure was four times higher than that in air at
800°C. However, the surface morphology showed the thickness of scale formed in sulfidizing gas mixture was still uniform, and the negligible concentrations of iron or sulfur together with relatively higher concentration of aluminium and oxygen indicated the formation of iron sulfide was suppressed by the growth of aluminium oxide layer. It was suggested that sulfur had still affected the processes controlling oxide scale growth without the formation of iron sulfides. The scale spallation after cooling to room temperature was also greater in the case of mixed gas exposure than in dry air.
Cyclic sulfidation experiments showed the protective scale could be quickly rehealed at 800°C. For iron aluminium alloys with Al concentration ranging from 16 to
22at%, the weight gain remained small if the Al content was higher than 18at%, and
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was dramatically increased when the Al content dropped down to 16at%. Both iron and aluminium sulfides were found in the scale formed on Fe-16Al.
All of the Fe3Al alloys containing 4-5%Cr showed higher weight gains that those alloys with 2%Cr. Small amounts of molybdenum(2%) or niobium were proved to be beneficial in decreasing the weight gain of Fe3Al containing 4% Cr, however the weight gain was still higher than alloys with lower Cr concentration. It remains uncertain whether the effect of chromium addition is to simply increase the sulphur diffusion by changing the defect concentration of an alumina scale or to cause mechanical defects to develop in the oxide scale through the growth of chromium sulfide nodules. Limited effects of Zr and Y on improving the scale adherence of Fe3Al in mixed gas were observed. In repeated thermal transients, no change in the scale appearance or composition resulted. Fe3Al alloys containing 4-5% Cr formed iron and chromium sulfides at the gas interface, and beneath this layer a mixture of Al2O3 and Al2S3 was detected. Initially formed scales fell off during cooling, however a slower scale formation rate were found after reheating to 800°C again. Therefore, less scale was lost during the subsequent cooling transient. Longer initial exposure resulted in a greater scale loss on the initial thermal transient and slower regrowth rate of the scale. Simply, this observation indicated the removal of the initially grown scale could decrease the corrosion rate in the subsequent stage. DeVan et al [186] suggested the possible explanation could be that sulfides formed in the initial stage were firmly attached the alloy surface when the samples were held at the high temperature. Furthermore, those sulfides limited the effectiveness of alumina as a barrier layer. After the spallation of the sulfides scale, alumina was the dominant scale constituent in the newly formed scale. The reason alumina was then the major scale constituent, but not in the initial scale, was thought to be the outwards diffusion of iron and chromium and enrichment of aluminium near the alloy surface during scale growth in the initial stage. 80 Literature Review
Klower [187] investigated iron aluminide containing aluminium concentrations varied from 6 to 17 wt%, and chromium concentrations between 2-10 wt% and alloy
800H (Fe-0.3Al-20Cr-31Ni) in different gas mixtures (Table 2.5 shows the details).
The results obtained are discussed below.
Table 2.5 Composition, Temperature range and Oxygen partial pressure of the test
gases [187]
Gas (in vol%) T(°C) Log PO2
Air 1100
CH4/H2: ac=0.8 850 -28
1000 -25
1100 -24
0.25% Cl2/ 20% O2/ Ar 850
1 and 5 and 10% SO2 in air 650-850
1% H2S/ 10% CO2/ H2 550 -26
Corrosion tests were cyclic. One cycle included 1.5 hours heating, 16 hours at reaction temperature and cooling down to room temperature. For oxidation, chlorination and carburisation tests, the entire testing time was 1008 hours (42 cycles), and 2016 hours (84 cycles) were used in the case of sulfidation experiments.
Weight gains of samples in dry air at 1100°C showed no large difference between different alloys, and their values ranges between 7 and 14 g/m2 after 42 cycles exposure.
Fe-10Al-2Cr and alloy 800H were tested in 1, 5 and 10% SO2/ air at 650°C, in
1%SO2/air at 750°C and in 1% SO2/ air at 850°C. At 650°C in 10% SO2/ air, the weight
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gain on the alloy Fe-10Al-2Cr was 3g/cm2, in contrast to the severe weight loss of alloy
800H caused by spallation. The corrosion rate of both alloys was increased with increasing SO2 concentration and temperature. In 1% H2S/1 0% CO2 / H2 gas mixture at 550°C, alloys with relative low Al concentration like Fe-5Al-2Cr, Fe-8Al-2Cr and
Fe-10Al-2Cr showed substantial weight loss. Other alloys like Fe-15Al-2Cr, Fe-17Al-
2Cr and Fe-16Al-10Cr showed only slightly weight gains after 2016 hours exposure. It was suggested by Klower that 12 wt% Al is needed to protect alloy under these corrosion condition.
The sulfidation/oxidation behaviour of iron aluminium bulk alloys was studied by
Kai and Huang [188]. Five Fe-Al binary alloys containing up to 40 at% Al were tested in H2/H2S/H2O mixture over the temperature range of 700°C -900°C. The partial pressures of sulfur and oxygen were 10-7-10-5atm and 10-24-10-20 atm, respectively. Iron based alloys containing 5, 10, 18, 28 and 40 at% Al were cast into ingots. Sulfidation experiments showed that the corrosion kinetics of the five Fe-Al alloys followed the parabolic rate law at all temperatures (See Figure 2.20 and Table 2.6). This indicated that the rate-controlling step was solid-state diffusion. In all cases, an initial transient stage followed by steady-state parabolic behaviour was observed. The duration of the transient stage decreased with increasing temperature and decreasing Al content. For iron aluminium alloys with Al content less than 18 at%, the reduction in corrosion rate was minor, a result similar to that of Klower [187]. Depending on the Al content and temperature, different scale microstructures and phase constitutions were obtained. For alloys with Al content less than 10 at% only sulfides were formed. Both sulfides and oxides were found on alloys contain 18-28 at% Al. At 700°C, only oxides were formed on Fe-40Al alloy, however, iron sulfide together with alumina were detected at temperatures higher than 800°C.
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Figure 2.20 Corrosion Kinetics for Fe-Al alloys: (a): Fe-5Al; (b) Fe-10Al; (c) Fe-
18Al; (d) Fe-28Al and (e) Fe-40Al [188].
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Table 2.6 Parabolic Rate constants (g2 cm-4 sec-1) and Apparent Activation Energies
of Fe-Al Alloys [188].
(a) in H2/H2S/H2O mixed-gases
Alloy (at%) 700°C 800°C 900°C Q (Kcal/ mole)
Pure Fe 5 95 x 10-8 1.88 x 10-7 3.64 x 10-7 20.6
Fe-5Al 9.14 x 10-9 5.06 x 10-7 2.44 x 10-7 37.1
Fe-10Al 5.68 x 10-9 3.72 x 10-8 2.13 x 10-7 40.9
Fe-18Al 1.80 x 10-9 1.91 x 10-8 1.27 x 10-7 48.3
Fe-28Al 2,46 x 10-11 5.70 x 10-9 1.74 x 10-8 75.5
Fe-40Al 6.59 x 10-13 1.87 x 10-11 1.02 x 10-9 82.6
(b) in pure sulfidation
2 -4 -1) Alloy (at%) Source Ps2 (atm) Temp. (°C) Kp(g cm sec Ref
-4 -9 Fe-9.8Al H2/H2S 10 700 3.83 x 10 [189]
-6 -7 Fe-9.0Al H2/H2S 10 900 8.90 x 10 [182]
-5 -7 Fe-18Al H2/H2S 10 900 2.90 x 10 [182]
-4 -7 Fe-18Al H2/H2S 10 900 4.90 x 10 [182]
-3 -7 Fe-18Al H2/H2S 10 900 4.90 x 10 [182]
By using XRD analysis, FeS and FeAl2S4 were found on low-Al alloys, FeS,
FeAl2S4 and Al2O3 were formed on Fe-18Al and Fe-28Al. In the case of Fe-40Al, only small amounts of FeS were detected. The cross section of the low-Al alloys showed
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duplex scales were formed, consisting of an outer layer of less-compact FeS and a heterophasic inner layer of FeS and FeAl2S4. The higher the Al content, the greater the amount of the FeAl2S4 was. Neither Al2S3 nor Al2O3 were formed on those low-Al alloys.
For Fe-18Al alloys after 8 hours exposure at 900°C, an inner mixture of Al2O3 and
FeAl2S4 and some localized corrosion products were observed. X-ray maps were used to reveal the scale constitution and showed aluminium were still remained in the inner
layer. Al 2O3, FeS and some FeAl 2S4 were found in the scale formed on Fe- 28Al alloys.
With increasing temperature, the amount of Al2O3 increased. Because of its ability to form a thin Al2O3 layer to protect the alloy, Fe-40Al formed many small rectangular nodule-shaped precipitations at 700°C, and result showed only a-Al2O3 formed. At
800°C, a small amount of FeS together with the major scale constituent Al2O3 formed on Fe-40Al. Examination of the surface morphology revealed that the shapes of Al2O3 included whiskers, small granules and nodules, while columnar FeS was formed on top of the Al2O3 at edges and the suspension hole. At 900°C, a layer of FeS formed on most area of the sample surface.
Marker studies were employed to analyse the rate-controlling step of Fe-Al alloys.
The thickness of the scale formed on Fe-28Al and Fe-40Al alloys was too thin to determine the Pt marker position. However, the Pt marker on Fe-18Al alloy corroded at
900°C revealed that the formation of the outer scale was controlled by outward diffusion of iron while the growth of the inner scale was due to inward diffusion of sulfur.
Recent research on the corrosion behaviour of Fe3Al intermetallic compound at
605°C and 800°C was undertaken by Lee and Lin [190], using 1%SO2/air gas mixture.
85 Literature Review
The thermogravimetric results for Fe3Al with 2%Cr in oxidation showed higher reaction rates than those in sulfidation/oxidation experiments at 605°C and 800°. It was suggested by these workers that SO2 might have a positive effect on the oxidation resistance of iron aluminide.
2.4.6.3.3 High temperature chlorination of iron aluminide alloys
Compared to the mechanisms of oxidation and sulfidation, chlorination is still a less well understood area [191]. Again, the potential of iron aluminides as materials with adequate corrosion resistance in chlorine-containing environments has been investigated
[129, 192-196].
Natesan [195] applied the electrospark-deposition (ESD) technique to coating type
316 stainless steel and alloy 800 substrates. These ESD coatings included Fe3Al with different bond coats of refractory metals and /or overlay coatings of noble metals. The corrosion tests were undertaken in Cl containing gas mixtures (p =1.2 x 10-23, 2 o2
p =5.2 x 10-10, p =9.4 x 10-17 and pHCl=2.1 x 10-3 atm) at 650°C for 1000 hours. s2 Cl2
The results are shown in Figure 2.21 [195]. It was found that the uncoated substrates were attacked by general corrosion and pitting, leading to weight losses unlike the relatively small weight gains of the coated alloys.
Three iron aluminide alloys together with one Fecralloy were tested in a simulated coal gasification atmosphere (CGA) containing HCl by Saunders et al. [193]. Table 2.7 shows the alloy compositions investigated in that research. Alloys were exposed in a
CGA gas mixture of composition 24.8%CO, 20.4%H2, 3.93%CO2, 6.3%H2O,
0.123%H2S balance N2. The additional HCl and H2O levels were controlled by passing the dry gas mixture over a heated HCl solution. Two HCl levels were chosen at
86 Literature Review
1000ppm and 3283ppm, respectively. The partial pressures of O2, S2 and Cl2 at the investigated conditions were also calculated and are shown in Table 2.8. At 450°C in
CGA + 1000 ppm HCl gas, similar weight changes were observed on Fecraloy, FAS1 and FA129 after 1000 hours exposure (0.3–0.8 g m-2). In contrast, FAPVIM was severely corroded after about 300 hours and its weight gain was about 10 mg cm-2.
From optical microscopy, it was shown that the material losses of FAPVIM alloy was about 50 mm. The scale thickness formed on the other alloys was much less than 1 mm.
At 550°C with HCl level at 1000 ppm, the mass changes of alloys fell into two groups again. As the oxygen partial pressure was also increased along with the increasing temperature, a protective oxide scale was formed on Fecraloy, FAS1 and FA129. Their weight gains ranged from 0.4 to 0.7 g m-2. Alloy FAPVIM showed 4.5 g m-2 after 1000 hours exposure, although it behaved well for the first 300 h of the experiment. In the gas mixture with higher HCl level (3283 ppm HCl), more severe corrosion was observed on Fecralloy and FAPVIM. Optical microscopy showed all alloys except
FA129 suffered severe corrosion attack. The ranking of alloy corrosion resistance of alloys in this case is: FA129 < FAS1 < Fecralloy < FAPVIM. However, localized attack was also found on FA129 and FAS1. The better corrosion resistance of FA129 and
FAS1 was attributed to their high aluminium concentration. However, the alloy with higher Cr content (FA129) showed better anti-corrosion performance. Since the CGA gas mixture pressure used in this study was much lower than that in most gasification plants, further investigation was suggested.
87 Literature Review
Figure 2.21 Weight changes for several iron aluminide coatings and uncoated
austenitic alloys exposed in gas mixtures containing H2S with and
without HCl [195].
Table 2.7 Alloy compositions in the study of Saunders et al. [193].
Alloy Composition (at%)
Designation Al Cr Fe Others
FAPVIM 16 5 Balance 1-Mo, 0.1Zr,
0.1Y
FA129 28 5 Balance 0.5-Nb, 0.2-C
FAS1 28 2 Balance 0.05-B
Fecralloy, 9.7 22.1 Balance 0.06-Y
SFEC*
· The nominal composition of Fecralloy in wt% is Al-5, Cr-22, Y-0.1, Fe-Balance
88 Literature Review
Table 2.8 Calculated partial pressures (Total pressure = 101325 Pa) [193].
Gas Species Partial Pressures Partial Pressures Partial Pressures
at 450°C, and at 550°C, and at 550°C, and
1000ppm HCl, Pa 1000ppm HCl, Pa 3283ppm HCl, Pa
-30 -26 -26 O2 4.13 x 10 1.12 x 10 1.12 x 10
-12 -11 -11 S2 4.24 x 10 4.00 x 10 4.00 x 10
-20 -18 -17 Cl2 9.32 x 10 1.99 x 10 2.15 x 10
In another study by Natesan et al [197], Fe3Al-based alloys containing alloying elements like Cr, B, Nb, C, Mo, and Zr, were exposed for about 48 h at 650°C in a
1vol% HCl-N2 gas mixture. Figure 2.22 shows the weight change of the alloys tested.
It was observed that the ternary and higher order alloys had smaller weight changes than those of Fe3Al. The critical concentration of Cr required to control the chlorination behaviour of Fe3Al-based alloy was concluded to be 5.5 wt%. Corrosion resistance and also the aqueous corrosion behaviour during shutdown (downtime corrosion, DTC) of iron aluminide coatings with surface Al contents of about 20-30wt% in simulated gasifier gases containing HCl has been investigated [196]. Table 2.9 and 2.10 show the experimental conditions and materials tested. Deposits (80% slagging gasifier fly ash,
10% carbon black, 5% NaCl and 5% FeCl3) based on those found in slagging gasifiers were coated on some of the samples. Pitting corrosion was found to be the main form of corrosion on aluminized T-11 steel (coated with a chloride rich deposit, but not exposed to DTC). Minor general and intergranular corrosion occurred as well. For
89 Literature Review
samples exposed to DTC, pits were found to have linked up and formed a continuous, but irregular corrosion front, and more than 60% coating was consumed. Alumina and iron oxide formed the main scale. Considerable amounts of FeS were detected as well.
The chloride content in the scale varied from 0.75 to 2.2 at %. It was explained that chlorine compounds in the gas attacked aluminium in the coating and formed volatile
AlCl3 which diffused out and reacted with H2O to form Al2O3 and HCl. DTC increased the concentration of chlorides at the scale metal-interface and thus accelerated the corrosion rate. FA129 was shown to have superior corrosion resistance in coal gasification environments no matter if the chloride rich deposit was present or not.
However, breakdown of the protective scale and massive pitting corrosion were observed after exposure to DTC conditions. The resulting irregular surface was attributed to overlapping pits. Once the protective Al2O3 scale was destroyed during
DTC, it could not be reformed, and the non-protective scale spalled off easily. SS310 steel developed a relatively thin Cr2O3 layer, which was quite protective, when the chlorine content was low (<1000ppm HCl). But a thick scale formed rapidly in the case of DTC conditions, and scale spallation occurred. Nonetheless, the overall corrosion rate of SS310 in DTC condition was still lower than those of iron aluminides since the protective scale formed on iron aluminides could not be reformed after spallation.
90 Literature Review
Figure 2.22 Weight Change data of Fe3Al-based alloys tested in a 1vol% HCl-N2
gas mixture at 650°C [197].
Table 2.9 Experimental conditions [196].
Pressure Gas composition, vol% Temperature Time HT DTC*
(atm) H2O, H2, CO, CO2, H2S, (°C) (hours) Cycles cycles
HCl
41 7.4, 43, 44, 5, 0.6, 0.04 540 572 3 2
1 3, 28.2, 64, 4, 0.8, 0.04 400 587 3 3
*Samples exposed to moisture during downtime (24 hrs, 30-35°C, H2O saturated air)
91 Literature Review
Table 2.10 Alloy composition (wt%) tested in simulated gasifier gases [196].
Al (wt%) Cr (wt%) Fe (wt%) Other
Aluminium diffusion coating 20-30 1-2 bal
FA 129 15.9 5.5 bal Nb 1,
SS310 tr 25 bal Ni 20,
Surface composition of Aluminized 20-25 1-1.24 bal Mo 0.5
T-11
92 Experimental Procedures
CHAPTER THREE
Experimental Procedures
3.1 SPECIMEN PREPARATION
A wide range of materials was chosen in the project, in the view of the complex experimental conditions. Materials compositions are shown in Table 3.1. Low carbon steel A1006 was used as the substrate for pack aluminizing process. Rectangular A1006 steel samples (10mm x 5mm x 1.5mm) were sectioned from 1.6mm thick steel sheet
(obtained from BHP Research). In the case of the aluminizing process, bare A1006 samples were ground on 600 grit SiC paper, ultrasonically cleaned in acetone (3mins) before using. The 2.25Cr1Mo steel was obtained from ANSTO. SG cast iron (SG 400) ingot was supplied by Carruthers Foundry Pty. Ltd. Zincalume ingot was supplied by
BHP Steel Pty. Ltd. Pure aluminium was obtained from Comalco Ltd. Materials investigated in corrosion tests were sectioned into 10mm x 10mm x 1mm samples, abraded with 1200 grit SiC paper and also ultrasonically cleaned in acetone (3 mins) prior to use.
In the case that NaF was used as the activator in the pack aluminizing process described in Chapter 4, bare A1006 samples were pre-coated with an yttria containing slurry before packing. The yttria slurry was prepared by mixing 4 mg pure yttria powder (particle size specification: 1-5mm; purity: 99.99%), 30ml acetone and 4ml organic binder (“Novaset 70/50” from Ashland Chemicals). Premixed yttria slurry was then applied to the specimen surface with an air-spray gun (Stan S-106, from Rich Star
Precision Industrial Co., Ltd.). The spray gun was operated at a pressure of 100 kPa.
93 Experimental Procedures Fe C Mn Si S Ni Cr Mo P Cu Al Others 1 A1006 Base .04 .26 <.004 .008 .02 .01 < .01 .015 .04 2.25Cr1Mo Base 0.15 03-0.6 0.5min 0.03 1.9-2.6 0.87- 0.03 1.13 SG cast iron Base ³2.5 £0.4 £2.5 £0.02 £0.08 (SG 400) 2 3CR12 Base .30- 1.25 1.00 ³.15 .50 12.00- .06 £.60 .40 14.00 253MA Base .10 0.80 1.4-2.0 .03 10.0- 20.0- N: 0.14- 12.0 22.0 0.20 Ce: 0.03- 0.08 SS304 Base .04- 2.0 1.00 0.03 8.00- 18.0- .10 11.0 20.0 3 Al 99.91 Zincalume 2 55 Zn: 43 4 Cu 99.99
Cu-S alloy
Table 3.1 Table of the composition of materials investigated (wt%)
94 Experimental Procedures
The thickness of the deposited yttria layer was controlled around 150 mm. Specimens coated with yttria slurry were then dried in air for 1 hour before packing. Figure .3.1 shows the cross-section of an un-coated steel sample with deposited yttria slurry.
3.2 PACK ALUMINIZING PROCESS
3.2.1 Pack Materials
Pure aluminium fine powder (particle size: ~60 mm; purity: 99.9% from
Goodfellow Cambridge Ltd.) was used as the master alloy in all the pack aluminizing processes described in Chapter 4. Either NH4Cl (Reagent grade; purity: 99.5% min.; obtained from AJAX Chemical Ltd.) or NaF (Reagent grade; purity: 99% min.; from
AJAX chemical Ltd.) was chosen as the activator. To prevent the master alloy from sintering, alumina powder (particle size: ~120mm; purity: 99% approximately; from
Ceramic Engineering Pty. Ltd.) was used as the inert filler material in all experiments.
Figure 3.1 Cross-section image of a specimen with pre-deposited Y2O3 slurry.
95 Experimental Procedures
3.2.2 Pack Preparation and Coating Deposition
Pack powder was premixed before being loaded into a cylindrical alumina crucible.
The proportions of pack constituents are listed in Table 3.2. Normally, 30g powder was mixed in every process. The powder mixture was put into a plastic container and then mixed for 0.5hr using a SPEX 8000 Mill/Mixer.
Table 3.2 Proportions of pack constituents used in pack aluminizing process
Proportion of pack constituents Case 1 Case 2
(wt%*)
5 part Master alloy Pure Aluminium Pure Aluminium
3 part Activator NH4Cl NaF
92 part Inert filler Al2O3 Al2O3
*Note: (unless otherwise stated, all compositions will be given in weight percent)
The powders and three A1006 steel specimens were then loaded into a cylindrical alumina crucible with dimensions of 30mm internal diameter and 50mm length approximately. The experimental apparatus and specimen arrangement inside the crucible are shown in Figure 3.2. The packed crucible was topped off with pack powder and covered with an alumina lid. Alumina based-cement was used to seal the lid.
Sealing cement was dried in air for at least 12 hr at room temperature.
The pack aluminizing process was carried out in a horizontal tube furnace. For each deposition temperature, the furnace temperature profile was measured to ensure the deposition process was taking place in the uniform hot zone. For a given aluminizing
96 Experimental Procedures
process, the packed crucible was placed into the uniform hot zone at room temperature.
High purity inert gas (either nitrogen or argon, from BOC Ltd.) at a flowrate of
100ml/min was introduced into the furnace at room temperature 20 mins after the furnace was sealed and throughout the whole process. To remove the residual moisture in the pack, the furnace was then heated up to 200°C for 2hr and held at this temperature for 2hr. The furnace was then heated up to the desired deposition temperature (900°C ) at a rate of 200°C/hr. Deposition times are reported as time at the depositing temperature. Therefore the heat-up and cool-down times were not counted in. Upon completion of deposition, the pack was allowed to cool in the furnace. Coated specimens were then removed from the pack and cleaned for examination.
Furnace
Ar out Ar in
Reaction tube (Alumina) Alumina crucible Thermocouple
e a-Bare A06 d-Lid b-Packing powder e-Alumina cement
c-Alumina crucible a b d c Figure 3.2 Schematic drawing of pack aluminizing process equipment.
3.2.3 Coating Analyses
Aluminized specimens were examined and characterized in the following aspects: weight gain, thickness, surface morphology and phase constitution.
97 Experimental Procedures
To analyse the aluminized specimens, X-ray diffraction (XRD, Siemens D5000) using Cu-Ka radiation was employed for surface phase characterization. One specimen was cross-sectioned and polished to a 0.25mm finish for metallography purpose with optical microscopy. FESEM (Field Emission Scanning Electron Microscope, Hitachi
5000) with EDAX (Energy Dispersive Analysis of X-Rays, operated at 20kv) was sometimes used for morphological studies. A Cameca SX50 electron probe micro- analyser (EMPA) was used for quantitative composition analysis using wave-length dispersive spectroscopy (WDS). All samples and EPMA standards were carbon-coated prior to analysis as some of the phases analysed were non-conductive (e.g.the oxide scale). The EPMA standards used were pure aluminium, iron, cooper, nickel, and
YAlO 3 (for yttrium and oxygen). An accelerating voltage and probe current of 15 keV and 20 nA, respectively, were used for all analysis. The intensity of the characteristic x- ray peak of each element was corrected by subtracting the background x-ray level and absorption and fluorescence was accounted for on an atomic number and sample composition basis via the usual matrix ZAF correction procedure [198].
3.3 HIGH TEMPERATURE OXIDATION
Aluminized A1006 specimens were then tested together with other materials in air at 800°C both isothermally and cyclically.
3.3.1 High Temperature Isothermal Oxidation
Materials with low melting point (such as zincalume and aluminium) and materials having unacceptable oxidation rate at the temperature investigated (e.g.: Cu-S alloy and copper) were not studied in the high temperature isothermal oxidation tests. High temperature isothermal oxidation experiments were carried out in either a TGA
98 Experimental Procedures
(thermogravimetric analyser, Cahn 2000) for kinetic studies or a horizontal furnace for long duration tests in air at 800°C.
The high temperature oxidation behaviour of A1006 steel, 2.25Cr1Mo steel, SG cast iron together with coated A1006 steel and 3Cr12 was determined using TGA.
Initial sample weights were measured before experiments started. Pure argon was used to purge the reaction tube for 20 minutes before air was led in. To eliminate the effect of gas flowrate on specimen, the same flowrate (100ml/min) was used for pure argon and air. After each exposure, a specimen was allowed to cool in flowing argon to room temperature. The specimen was then removed and its final weight recorded. A schematic drawing of the TGA apparatus is shown in Figure 3.3.
Only coated A1006 steel, 304 grade stainless steel and 253MA were tested in the horizontal furnace shown in Figure 3.4. Their initial weights were measured before they were loaded in an alumina boat. Specimens were located in the cold zone of the furnace (at 80°C approximately) initially and then pushed into the hot zone gradually with a steel rod. Pure nitrogen with a flowrate of 100ml/min was used to purge the furnace before the specimen reached the reaction temperature (normally 5 mins after the alumina boat was settled in hot zone). Air (flowrate of 100ml/min) was used as the reaction gas. Upon completion of the experiment, air was turned off and the alumina boat was pulled out gradually with the protection of pure nitrogen again. Specimens were removed and then weighed.
Corrosion products were analysed using XRD. Specimens were then mounted in resin, cross-sectioned and polished for metallography purpose. FESEM and FIB
(Focused Ion Beam Miller) were sometimes used to study the surface morphology.
99 Experimental Procedures
TGA balance
Gas exit Gas entry Computer
Furnace controller Silica reaction tube
Furnace
Sample
Figure 3.3 Schematic drawing of thermogravimetric analyser
Furnace
Gas in Gas Out
Alumina boat
Figure 3.4 Schematic drawing of the high temperature apparatus.
100 Experimental Procedures
3.3.2 High Temperature Cyclic Oxidation
High temperature cyclic oxidation tests were carried out on some materials, including 304 grade stainless steel, 253MA, coated and un-coated A1006 steel using a vertical furnace shown in Figure 3.5. Each cycle included 1 hour at reaction temperature 800°C and, 15 mins cooling in laboratory atmosphere.
The weight changes of specimen were measured after designated numbers of cycles. XRD was used to analyse surface phase constitution. Specimens were then cross-sectioned and mounted for further metallographic study.
3.4 CORROSION TESTS IN MIXED GASES
Materials were investigated in different gas mixtures at 400°C and 800°C, respectively. These corrosion tests were also undertaken using a horizontal furnace.
Specimens were loaded in an alumina boat, and the procedures were similar to those used in high temperature oxidation experiments.
3.4.1 Corrosion Tests in Gas Mixtures at Low Temperature
All materials were investigated in four different gas mixtures in order to determine the effects of different gas species on the corrosion behaviour of materials. The compositions of these gas mixtures are shown in Table 3.3. In the case that hydrogen chloride was used, air (or O2 + N2) was passed through water (thermostat temperature:
27°C) or 18wt% HCl solution (thermostat temperature: 37.5°C) to obtain the required
HCl and water vapour partial pressures [199]. A schematic drawing of the apparatus is shown in Figure 3.6. After the completion of experiments, specimens were removed and their weight changes recorded. XRD was used to analyse corrosion products.
101 Experimental Procedures
Optical metallography was carried out and FESEM with EDAX was sometimes used in analysing both surface morphology and cross-sections.
Specimen
Vertical Furnace Control Box
Figure 3.5 Schematic drawing of cyclic oxidation apparatus
3.4.2 Corrosion Tests in Gas Mixture at High Temperature
As described in the high temperature oxidation section, zincalume, aluminium, copper and Cu-S alloy were not tested at this temperature. The details of the experimental procedure were similar to those of the high temperature oxidation experiments. Specimen weight gain, thickness, surface morphology and phase constitution were studied using XRD, optical microscopy and FESEM with EDAX were used.
102 Experimental Procedures
Table 3.3 Experimental conditions for mixed gas corrosion.
400°C 800°C
Ø Gas 1: Isothermal oxidation in air Ø Isothermal sulfidation in 24.9%
-3 Ø Gas 2: Air + H2O (0.0375atm) CO + 30.2% CO2 + 4.98 x 10
-4 Ø Gas 3: Air + HCl (5´ 10 atm) + H2O atm SO2 + N2 (balance)
(0.0375 atm)
-4 Ø Gas 4: Air + HCl (5´ 10 atm) + SO2
-4 (8.5x 10 atm) +H2O (0.0375 atm)
Flowmeter Toggle valve Toggle valve 3 ways valve O2/Air N2 18wt% HCl solution SO2/N2 3 ways valve
Thermostat
Samples Furnace
H2O NaOH
Figure 3.6 Schematic drawing of the apparatus used for low temperature corrosion tests.
103 Pack Cementation Process
CHAPTER FOUR
Pack Cementation Process
4.1 INTRODUCTION
An aluminium-rich coating was developed on A1006 low carbon steel using the pack aluminizing process. In this chapter, coating growth kinetics and morphological development are reported.
4.2 THERMODYNAMIC CALCULATION
The efficiency of NH4Cl and NaF activator has already been described in chapter
2.4.6.2 [141, 148]. To establish a suitable pack aluminizing system, a number of free energy minimization calculations (similar to those in the study of Rapp et al [142] and
Da Costa [200]) were carried out for different activators.
The computer program CHEMIX was used to calculated the equilibrium conditions
(including partial pressures and composition) prevailing in the bulk pack at the aluminizing process temperature. CHEMIX is a part of CSIRO-SGTE
THERMODATA system [201] and utilizes a computational procedure based on free energy minimization [152] (Appendix 1 shows the input free energy data of those calculations). Partial pressures were calculated for packs at their initial composition under an inert (argon) atmosphere and at 900°C. Pure iron was used as the substrate specification.
The input values for pack materials in the calculation were based on 23.4g packs comprised of 3wt% activator (NH4Cl or NaF), 5wt% master alloy (pure aluminium and
104 Pack Cementation Process
92wt% a-Al2O3. The amount of iron used in the calculation was taken to be 3g in total.
The input value for argon was assumed to be the amount which occupied the free pack volume (assuming the pack to have 45% porosity) at the processing temperature. Ideal gas behaviour was. The equilibrium bulk-pack partial pressures and those condensed phases present for packs with different activators at 900°C are listed in Table 4.1.
Data for the free energy of formation for all the gaseous and condensed species were calculated using thermodynamics calculation software – THERMOCALC [202].
According to the activity data shown in Table 2.3 (chapter 2.4.4), iron aluminides with up to 75at% aluminium are possible aluminizing products in both packs using pure aluminium masteralloy (aAl=1).
4.3 RESULTS
Deposition times are recorded from the time the furnace reached the deposition temperature. Thus, for all of the results present below, zero hour is the time at which the furnace reached deposition temperature unless otherwise stated.
4.3.1 Coatings Formed on Steel Using Different Activators without Yttria
Deposition
Bare A1006 steel specimens were coated at 900°C for 8hrs using either NH4Cl or
NaF activator. Images of the cross-sections of coatings formed using different activators are shown in Figure 4.1. XRD results for coatings formed using different activators are shown in Figure 4.2.
As shown in Figure 4.1a, a total thickness about 32mm coating was formed using
NH4Cl activator. Large amounts of second “phase” inclusions were seen at the dark regions in the coating surface down to the inside of the coating. EDS results showed
105 Pack Cementation Process
that those black inclusions inside the coating were rich in aluminium and oxygen, whereas those close to the coating surface were cavities. It is believed that the black phase is alumina from the pack materials. Aluminium concentrations within the coating are shown in Figure 4.3. XRD results showed that only FeAl and Fe3Al could be detected, and this was consistent with the aluminium concentration results measured using EDS. As these alumina inclusions inside the coating can be considered as markers, the coating formed on A1006 steel using NH4Cl was mostly controlled by the outwards diffusion of iron. Of course, the inwards diffusion of aluminium also occurred throughout the process.
A more promising coating with only a small amount of inclusions was formed using NaF activator. From the optical image of its cross-section, the relatively thick coating (total thickness of 90mm) can be categorized as a multiple layer coating. It consisted of an outer bright coloured layer (thickness » 30-40mm), a thin intermediate grey layer and a light grey colour inner layer with fine inclusions. Iron aluminides formed on the surface were analysed using XRD, and results are shown in Figure 4.2.
Compared with those phases formed on A1006 steel using NH4Cl activator, the assemblage iron aluminides formed using NaF activator contained more aluminium.
Black inclusions were analysed using EPMA and again were found to be rich in aluminium and oxygen. Therefore, it can be concluded that the coating formed on
A1006 steel using NaF activator is also produced by the outwards diffusion of iron.
Elemental concentrations were studied using EPMA, and results are shown in Figure
4.4. The thickness of coating measured by EPMA is slightly thicker than the previous measurement using optical method. Coating phase constitution was also defined, and the results showed to be very similar to those XRD results.
106 Pack Cementation Process
Table 4.1 Partial pressures of species and condensed phases in packs using
different activators.
NH4Cl activator NaF activator
Species partial pressures Al 1.26E-08 Al 1.79E-08
(atm) AlCl 2.42E-02 AlF 5.56E-03
AlCl 2 1.67E- 04 AlF 2 7.56E- 05
AlCl 3 0.1066 AlF 3 1.96E- 04
Al 2Cl 6 1.60E- 04 Al2F6 7.48E- 07
Cl 2 7.2E- 17 Na 6.31E- 03
AlN 1.5E-15 NaF 2.5E-03
NH3 5.307E-05 Na2F2 6.81E-07
HCl 2.4E-04
Condensed phase Fe (s) Fe (s)
Al2O3 (s) Al2O3 (s)
Al (l) Al (l)
Na F (s)
107 Pack Cementation Process
(a)
(b)
Figure 4.1 (a) SEM image of the cross-section of coating formed using
NH4Cl activator
(b) Optical image of the cross-section of coating formed using
NaF activator (etched by 5% Nital)
108 Pack Cementation Process
(a)
(b)
Figure 4.2 (a) XRD scan result for coating formed using NH4Cl activator
(b) XRD scan result for coating formed using NaF activator
109 Pack Cementation Process
35
30
25
20
15
10
5
Aluminium Concentration (at%) 0 0 5 10 15 20 25 Thickness (mm)
Figure 4.3 Aluminium concentration of coating formed using NH4Cl activator at
900°C after 8hr.
4.3.2 Coating formed using NaF activator on steel pre-coated with yttria slurry
According to the observed superior performance of NaF, this activator was used for all further studies. Yttria slurry was also applied on the un-coated A1006 steel surface.
Yttria slurry acts as a physical barrier between substrate and alumina. Figure 4.5 shows an optical image of a cross-section of the coating formed on A1006 steel (pre-coated with yttria slurry) using NaF activator. As seen in the optical image, coating formed can be described as a combination of a thin outermost layer of non-uniform thickness above a relatively thick inner layer containing a small amount of black inclusions. A branch-like phase can be seen at the inner front of the coating. Using HF-containing etch solution to distinguish some iron aluminides having different aluminium level by showing different colour, phase with a dark grey colour was found to be FeAl, while
110 Pack Cementation Process
Fe3Al showed a light grey colour. Some micro-cracks can also been seen to cross
through the outermost layer and part of the inner layer. XRD results for the surface
phases are shown in Table 4.2. Except for some aluminium yttrium oxides and yttrium
oxides, there was no significant difference in the iron aluminides compared to those
found in the case where yttria slurry was not applied. EPMA analysis showed that the
black inclusions inside the coating were yttrium-containing oxides. Clearly, a relatively
pack inclusion free coating was produced after applying yttria slurry.
120
FeAl Fe Al a-Fe 100 2 FeAl 3
80 FeAl2+FeAl
Fe 60 Al
40
20 Elements concentration profile (at%)
0 0 50 100 150 Thickness (mm)
Figure 4.4 EPMA result for coating formed using NaF activator after 8hr at 900°C .
111 Pack Cementation Process
500 x
Y2 O3
10 m m
Figure 4.5 Cross-sectional optical image of the of coating (etched by 1 part of HF
+ 2 parts of Nital + 1 part of H2O) formed on A1006 steel (pre-coated
with yttria slurry) at 900°C after 8hr.
Careful EPMA analysis of the coating was carried out to study its phase constitution. Figure 4.6 shows the EPMA result of coating (with pre-deposited yttria slurry) formed using NaF activator at 900°C after 8hr aluminizing treatment. The coating thickness measured using EPMA was about 160mm, which was 20mm thicker than that formed on A1006 steel without an yttria slurry deposit. Although XRD detected aluminium yttrium oxides on the coating surface, no significant amount of yttrium could be found here by EPMA analysis. However, yttrium and oxygen were detected at measurable levers 50mm beneath the coating surface. This is another evidence that coating formed on A1006 steel (pre-deposited with yttria slurry) is controlled by the outwards diffusion of iron.
112 Pack Cementation Process
coated) at 900 ° C - . EPMA results for coating formed on A1006 steel (yttria slurry pre 4.6 Figure after 8hr
113 Pack Cementation Process
4.3.3 Surface morphologies of coating grown using NaF activator
To understand the process of coating formation, a surface morphology study was undertaken on coatings formed on A1006 steel after different exposure times. FESEM images of the coating surfaces shown in Figure 4.7.
As seen in Figure 4.7, some coating had already been formed at “zero time” during the heating up period. A net-like structure can be observed, and it seems that the substrate grain boundary was the favoured location for coating to grow. Scratches from grinding could also been seen. Large numbers of cavities could also be found on the surface and it is believed that they were left by pack materials. After 2 hr aluminizing treatment, the coating started to grow more generally on the substrate, and after 4hr, a uniform coating had already been built up. The grain boundaries still can be seen.
Using EDS analysis, the white particles observed on the surface were found to be some pack materials. There were no obvious differences between SEM images of coatings formed after 6hr and 8hr.
4.3.4 Phase constitution of coatings formed for various deposition times
Cross-sectional optical images of coatings formed after different deposition times are shown in Figure 4.8. Their XRD analysis results are shown in Figure 4.9. For the coating formed at zero time, a thin layer of grey colour phase can be seen. The front of the grey phase was found to grow much faster along the grain boundary. A coating formed after 2hr has already formed multiple layer structure similar to that formed after
8hr, but much thinner. With increasing deposition time, the thickness of the top bright phase (iron aluminides with higher aluminium levels) increased. As seen in Figure
4.8d, the coating sometimes contained more yttria inclusions.
114 Pack Cementation Process Figu re 4.7 Surface morphology of coating formed after different deposition time.
115 Pack Cementation Process
Table 4.2 XRD results for coating formed on A1006 steel (pre-coated with
yttria slurry) at 900°C after 8hr.
XRD Result
Iron Aluminides Fe2Al5
FeAl 2
FeAl
Fe3Al
Aluminium yttrium oxides Al5Y3O12
AlYO3 Al Y O 2 4 9
Yttrium oxides Y2O3 YO1.401 YO1.458 YO1.335
116 Pack Cementation Process
(a)
(b)
117 Pack Cementation Process
(c)
(d)
118 Pack Cementation Process
(e)
Figure 4.8 Optical images of the cross-sections of coatings (etched)
formed for various deposition times at 900°C
(a) 0hr (b) 2hr (c) 4hr
(e) 6hr (e) 8hr
119 Pack Cementation Process Figure 4.9 XRD results of coatings formed for various deposition times.
120 Pack Cementation Process
EPMA was used to determine the phase constitution of coatings formed after different treatment time. Figure 4.10 shows the elemental concentration profiles.
120
100
80 Al Fe 60 Y O 40
20
Elements concentration profile (at%) 0 0 5 10 15
Thickness (mm)
(a)
120
100
80 Al Fe 60 Y O 40
20
Elements concentration profile (at%) 0 0 20 40 60 80
Thickness (mm) (b)
121 Pack Cementation Process
120
100
80 Al Fe 60 Y O 40
20 Elements concentration profile (at%) 0 0 20 40 60 80 100 120 Thickness (mm)
(c)
120
100
80 Al Fe 60 Y O 40
20 Elements concentration profile (at%) 0 0 20 40 60 80 100 120 140 Thickness (mm)
(d)
122 Pack Cementation Process
120
100
80 Al 60 Fe Y 40 O
20
Element concentration profile (at%) 0 0 50 100 150 200 Thickness (mm)
(e)
Figure 4.10 EPMA results for coatings formed for various deposition times
(a) 0hr (b) 2hr
(c) 4hr (d) 6hr
(e) 8hr
4.3.5 Coating growth kinetics
The growth kinetics of coatings formed in NaF activator pack at 900°C are presented in the form of a parabolic plot in Figure 4.11b. Thickness data used in the
123 Pack Cementation Process
plot are obtained from EPMA results and shown in Figure 4.11a. However, an error existed during coating thickness measurement since the thickness varied with samples location. From Figure 4.11, it is seen that the coating formation rate is approximately parabolic [203]: 2 X =kpt (4.1) where X is the coating thickness, kp is the parabolic rate constant and t is the deposition time, indicating that the coating growth was diffusion controlled. The value of kp was
-9 2 -1 calculated to be 3.5x10 cm sec , consistent with solid-state diffusion.
4.4 DISCUSSION
4.4.1 Effect of activator type on coating formed on A1006 steel
As described in section 4.3.1, coatings formed using different activators are different in their microstructures, coating thickness and phase constitution. As the other important parameters in coating process (deposition time, temperature, ratio between masteralloy and activator) were fixed, it is considered that the activator type was the main reason causing those differences.
According to the data for aluminium activities of iron aluminides at 900°C [185] listed in Table 2.3, Fe2Al5, FeAl2, FeAl and Fe3Al should be formed on A1006 steel using either activator. However, the XRD results for coating formed using NH4Cl revealed FeAl and Fe3Al only. When NaF is the activator, all iron aluminides expected were found. Assuming that the equilibrium between pack/metal was established in the initial stage, all expected iron aluminides would have formed on A1006 steel using
NH4Cl activator as well. However, interdiffusion of iron and aluminium would then have converted the higher aluminides to FeAl and Fe3Al if the supply of gaseous aluminium bearing species was insufficient to maintain the surface aluminium activity
124 Pack Cementation Process
As pointed out by other workers [141, 148], the efficiency of volatile activators
(e.g. NH4Cl), depends more on the initial amount of activator used than the temperature, whereas for stable activators, the efficiency of the system largely depends on the temperature. As described in section 2.4.6.1, aluminium will react with HCl which derives from decomposition of NH4Cl (Eq. 4.2- Eq. 4.6).
NH4Cl(g)=NH3(g)+HCl(g) (4.2)
2NH3(g)=N2(g)+3H2(g) (4.3)
Al+HCl(g)=AlCl(g)+1/2H2(g) (4.4)
Al+2HCl(g)=AlCl2 (g)+H2(g) (4.5)
Al+3HCl(g)=AlCl3(g)+3/2H2(g) (4.6)
In the present study, because the alumina crucible was in semi-sealed condition, it was possible for these HCl to escape from the crucible. In this event, the supply of aluminium-bearing volatile species would be curtailed, leading to a decreased surface aluminium activity and destabilisation of the higher aluminides. The thinner coating produced by NH4Cl activated packs is also thereby explained.
4.4.2 Surface morphology and kinetics of coating formation
From both surface morphology and cross-section images (Figure 4.7 and Figure
4.8) of coatings formed for various deposition times, it was found that the deposition of aluminium on substrate had already started during the heating up period before “zero” time. These net-like microstructures found on the coating surface at “zero” time indicated that diffusion along the substrate grain boundary is the dominant diffusion type initially. It seems likely also that nucleation of aluminide is favoured at grain
125 Pack Cementation Process
200 180 160 140 m) m 120 100 80
Th ickness ( 60 40 20 0 0 2 4 6 8 10 Time (hours)
(a)
35
30 ) 2 25 x cm
-5 20 (10 2 15
10 Th ickness 5
0 0 5 10 15 20 25 30 35
Time (103 x sec)
(b)
Figure 4.11 (a) Diagram of coating thickness vs. time (using NaF activator) at
900°C
(b) Diagram of coating thickness2 vs. time (using NaF activator) at
900°C
126 Pack Cementation Process
boundaries. These light grey colour branch tips shown in the cross-sectional images
(Figure 4.8b-e) indicate that diffusion along the grain boundary is also the dominant diffusion type during cooling down period. During coating growth, intragranular diffusion became important, leading to coatings of uniform thickness. The coating formation mechanism can be defined as a combination of outwards diffusion of iron and inwards diffusion of aluminium. This was demonstrated by the location of inert markers at intermediate position within the coatings. The Kinetics of coating formation were parabolic indicating that solid-state diffusion is the controlling-step.
Micro-cracks were found in the cross-sectional images of coating formed after 4hr deposition. This can be attributed to the multiple-layer coatings formed. Different iron aluminides have their own different thermal expansion coefficients. In the case that multiple-layer coatings were built up on the substrate, stress could be created between different iron aluminides. During the cooling down period, not all the stress can be relieved. Therefore, those residue stresses were relieved through the formation of micro-cracks.
The kinetics of coating formation are parabolic as the result that solid-state diffusion is the controlling-step.
4.4.3 Effect of yttria slurry deposit on coating constitution
To obtain the pack inclusion free aluminized coating, yttria slurry was used as a physical barrier between pack and coating. XRD results obtained for coatings with pre- deposited yttria slurry indicated that aluminium yttrium oxides were formed, although they were not found using FESEM with EDS. Instead of having alumina inclusions, small amount of yttria were identified inside the coating using EPMA.
127 Pack Cementation Process
4.5 CONCLUSIONS
i. An efficient pack mixture has been developed to obtain aluminium-rich coating
on A1006 steel. The mixture is based on use of pure aluminium master alloy at
a 5%level, and 3% NaF activator. The deposition temperature was 900°C.
ii. The coating formed on A1006 steel after 8hr can be described as a multiple-layer
coating, consisting of an outmost layer of iron aluminides having high
aluminium level (Fe2Al5, FeAl2), an intermediate layer of FeAl and an inner
layer of Fe3Al.
iii. With using pure aluminium master alloy, a coating formed by both outwards
diffusion of iron and inwards diffusion of aluminium was obtained. The
kinetics of the coating formation were parabolic, and kp rate was consistent with
solid-state diffusion control.
iv. Application of an yttria slurry deposit provided a physical barrier between pack
and coating. As a result, the coating formed with yttria slurry deposit was
essentially free of pack inclusions.
128 High Temperature Oxidation
CHAPTER FIVE
High Temperature Oxidation
5.1 INTRODUCTION
The production of aluminized coating on A1006 steel (pre-deposited with yttria slurry) using NaF activator was described in Chapter 4. To evaluate its high temperature oxidation resistances, aluminized A1006 steel, together with some other competitive materials, was tested at 800°C in air. Due to the different service conditions that may apply in industry, materials were tested both isothermally and cyclically.
5.2 RESULTS
5.2.1 Isothermal Oxidation
As described in the literature review, iron aluminides have superior high temperature oxidation resistance. To assess the high temperature oxidation performance of aluminized steel, isothermal oxidation experiments were carried out using TGA and horizontal furnace.
5.2.1.1 TGA Experiment
Coated A1006 steel, bare A1006 steel, 2.25Cr1Mo steel, SG cast iron, 3Cr12 steel,
253MA and SS304 stainless steel were exposed to air in a TGA apparatus at 800°C for various times Experimental details were provided in Section 3.3. The oxidation kinetics
129 High Temperature Oxidation
measured in this way are shown in Figure 5.1. Oxidation rules were rapid for uncoated
A1006, and 2.25Cr1Mo steels, and moderate for the other materials.
As expected, thick grey coloured scales were formed on A1006 steel and
2.25Cr1Mo steel. For SG cast iron, a relative thin scale, with small black nodules was observed. No scale of significant thickness could be found on coated A1006 steel,
3Cr12, SS304 or 253MA steels. However, the surface colour changes observed on these chromium-containing steels after exposure indicated a thin film of oxide layer has developed on these steels. The corrosion products formed on these materials were then analysed using XRD. Table 5.1 shows their XRD results.
Reacted materials were then cross-sectioned. Metallographic images for A1006 steel, 2.25Cr1Mo and SG cast iron are shown in Figure 5.2. For A1006 steel, a thick two-layered scale was formed (Figure 5.2a). Cracks through the whole scale thickness and gaps between scale and base metal can also be seen. The 2.25Cr1Mo steel formed a thinner scale than that formed on A1006 steel. Underneath the external scale, an internal precipitate zone also developed. A thin, two layered scale was formed on SG cast iron. In the region where graphite had been exposed to air, the scale was found to grow underneath the exposed spherical graphite.
FESEM was employed to study the surface morphologies of the oxides grown on
3Cr12, SS304 and 253MA steel, and images are shown in Figure 5.3. It is seen that local attack had occurred on all three materials. The attack was more general and severe on 3Cr12 than on the other two materials. Scratches left from grinding can still be seen.
A flake-like oxidation product was found. Analysis using EDS showed that the oxidation product was rich in chromium and oxygen, coincident with the XRD results
130 High Temperature Oxidation
131 High Temperature Oxidation
132 High Temperature Oxidation
133 High Temperature Oxidation
134 High Temperature Oxidation
of Table 5.1. Surface morphology studies of the oxide layer formed on coated A1006 steel was carried out only on samples in the later long-term oxidation experiments.
For coated A1006 steel, 3Cr12, SS304 and 253MA steel, no visible scale could be found by cross-sectional metallographic examination.
Table 5.1 XRD results for materials oxidized in air at 800°C in TGA experiments
Materials XRD results
A1006 steel Wustite, Hematite, Magnetite
2.25Cr1Mo steel Wustite, Hematite, Magnetite
SG cast iron Wustite, Hematite, Magnetite
3Cr12 steel Cr2O3
SS304 steel Cr2O3, FeCr2O4
253MA Cr2O3
Coated A1006 steel g,q, a-Al2O3
135 High Temperature Oxidation
(a) A1006 steel
(b) 2.25Cr1Mo steel
136 High Temperature Oxidation
(c) SG cast iron
Figure 5.2 Cross-sectional images of A1006 steel, 2.25Cr1Mo steel and SG cast
iron after oxidized (TGA) at 800°C .
137 High Temperature Oxidation
3Cr12
15 mm
(a)
3Cr12
1 mm
(b)
253MA
15 mm
(c)
138 High Temperature Oxidation
SS 304
1.5 mm
(d)
Figure 5.3 SEM images for 3Cr12, SS304 and 253MA oxidized in air at 800°C in
TGA experiments.
139 High Temperature Oxidation
5.2.1.2 Long Time Oxidation Experiments
The TGA experiments showed that some materials had better oxidation resistances: coated A1006 steel, SS304 and 253MA steel. Long term oxidation experiments for these materials were then carried out in a horizontal furnace at 800°C for different exposure times.
The weight gains of these tested materials were measured after each exposure and results are shown in Figure 5.4. After exposure in air for 2500 hours at 800°C, there were still no significant amounts of oxidation product built up on these samples. The
253MA steel showed weight loss due to partial scale spallation during the cooling down period at the end of each exposure. In the case of SS304 steel, partial scale spallation may have affected observed weight gains at longer exposure time.
The oxidation kinetics for SS304 stainless steel and coated A06 are shown in
Figure 5.5 to be parabolic. The oxidation kinetics for 253MA steel could not be calculated due to its partial scale spallation. Because no weight change was found for coated A1006 steel between 1680 and 2500h, this data was ignored when calculating its parabolic rate. The parabolic rate constants are shown in Table 5.2.
Table 5.2 Parabolic rate constants for oxidation in air at 800°C
2 -4 -1 Material Parabolic rate (kp: g .cm .sec )
Coated A1006 steel 3E-14
SS304 stainless steel 2E-14
140 High Temperature Oxidation
of coated A1006, SS304 and 253MA steel oxidized in Weight changes
Figure 5.4 air at 800 ° C for different exposure times.
141 High Temperature Oxidation
.
Figure 5.5 Oxidation kinetics of coated A1006 and SS304 steel in air at 800 °C
142 High Temperature Oxidation
The oxidation products formed on these three materials after each trial were then analysed using XRD, and the results are summarised in Table 5.3. The detailed results are shown in Figure A.3 (Appendix 2). For SS304 and 253MA steel, their XRD results showed strong peaks corresponding to the base metal, as a result of the very thin oxide layer(s) formed, or the partial scale spallation. The XRD results for coated A1006 steel are relatively complex. In addition to phases in the metal base (FeAl, FeAl2, Fe3Al,
Al5Y3O12, AlYO3, Al2Y4O9 and yttrium oxide(s)), metastable alumina phases and stable a-Al2O3 can also be detected. In Figure A.3 g-i, only oxide phases are labelled.
Table 5.3 XRD results for oxide scales grown in air at 800°C
Materials XRD results
Coated A1006 steel g, q, b and a phase Al2O3
SS304 steel FeCrO4, Cr2O3
253MA steel FeCrO4, Cr2O3
FESEM was used to study the surface morphology of these three materials oxidized for 2500 hours, and representative images are shown in Figure 5.6 and Figure 5.8, respectively.
As shown in Figure 5.6a, a thin, plate-like scale was formed on coated A1006 steel after long term oxidation. A ridged structure was found to be formed on the top of the plate-like scale. Yttrium-containing oxidation product was occasionally found on the steel surface and is shown in Figure 5.6b. In addition to these structures, needle-like
143 High Temperature Oxidation
15 m m
Figure 5.6a
Yttrium rich particle
m 6.67 m
Figure 5.6b
Figure 5.6c
144 High Temperature Oxidation
Sub-layer
6 mm
Figure 5.5d
Figure 5.6 SEM images for coated A1006 steel oxidized in air at 800°C for 2500
hours.
145 High Temperature Oxidation
oxide growth could also be found in some areas (Figure 5.6c). As seen in Figure 5.6d, another oxide layer (sub-layer), with the same structure as the scale, had formed underneath a damaged scale. Examination by EDS showed all oxidation products to be rich in aluminium and oxygen.
As seen in Figure 5.6, a relatively compact aluminium oxide layer formed on coated A1006 steel during oxidation. However, it is in principle for iron oxide(s) to grow in some areas where aluminium is consumed. Aluminium can be consumed either by the inwards diffusion of aluminium from coating into the steel or by the formation of aluminium oxide(s). In the latter case, the consumption of aluminium can be accelerated by the continuing scale spallation. It is therefore important to study the re- healing ability of the coating. FIB was used to achieve this purpose. Figure 5.7 shows images of sections milled through oxidized coating surface regions. It is seen that the compact oxide layer formed on coated A1006 steel was only about 1mm thick. No evidence of the formation of iron oxide(s) could be found. Underneath a through-scale crack (Figure 5.7b), a sub-layer with the same structure as the outermost scale had been reformed to provide further protection to coated A1006 steel.
In Figure 5.8, similar surface morphology to those observed in TGA experiments was found. Cracks along the grain boundaries and network-like oxide(s) can be seen on both materials. EDS result shows those network-like oxidation products are rich in chromium and oxygen.
5.2.2 HIGH TEMPERATURE CYCLIC OXIDATION
Bare A1006 steel, SS304, 253MA and coated A1006 steel were tested for high temperature cyclic oxidation resistances, using the techniques described in Section
3.3.2. Their weight changes were shown in Figure 5.9, where extensive reaction of bare
146 High Temperature Oxidation
2 mm
(a)
2 mm
(b)
Figure 5.7 SEM images for coated A1006 steel oxidized in air at 800°C for 2500
hours.
147 High Temperature Oxidation
(a)
(b)
(c)
Figure 5.8 SEM images for SS304 (a) and 253MA (b, c) oxidized in air at 800°C
for 2500 hours.
148 High Temperature Oxidation
A1006 steel is obvious, but little net weight change for coated A1006 or the heat resisted steels.
As expected, a thick black scale was formed on bare A1006 steel, whereas thin grey scales were found on SS304 and 253MA steels. Under cyclic oxidation condition,
SS304 showed weight loss due scale spallation. After 200 cycles, small weight gain was obtained on coated A1006 steel and 253MA steel. Partial scale spallation was also occurred on 253MA during oxidation.
The oxidation products formed on these materials were then examined using XRD, and the results are shown in Table 5.4.
Table 5.4 XRD results for materials oxidized in cyclic oxidation experiments
Materials XRD results
Bare A1006 steel Wustite, Hematite, Magnetite
Coated A1006 steel a-Al2O3
SS304 steel Cr2O3, FeCr2O4
253MA steel Cr2O3
Reacted samples were then cross-sectioned, and metallographic images are shown in Figure 5.10. A two-layered scale consisted of a light grey colour outmost layer and a thick grey colour inner layer formed on bare A1006 steel. The thinned thickness of the remaining steel indicates the massive consume of iron. In the case of 253MA steel, a
149 High Temperature Oxidation
140 A06 120
100 ) -1 80
60 W/A (mg cm D 40
20
0 0 25 50 75 No. of cycles
5
0 0 50 100 150 200 250 -5 ) -1 -10
-15 W/A (mg cm D -20 Coated A06 SS304 -25 253MA -30 No. of cycles
Figure 5.9 Weight changes of bare A1006 steel, SS304, 253MA and coated A1006
steel in cyclic oxidation experiments.
150 High Temperature Oxidation
A06 50 cycles
(a)
253MA 200 cycles
(b)
151 High Temperature Oxidation
304 200 cycles
cavities
(c)
Coated A06 200 cycles
(d)
Figure 5.10 Cross-sectional images for bare A1006 steel, 253MA, SS304 and
coated A1006 steel oxidized in cyclic oxidation experiments.
152 High Temperature Oxidation
thin scale of dark grey colour can be seen. However, the scale shown on Figure 5.10b is discontinuous and had perhaps been damaged by partial scale spallation. A duplex scale formed on SS304 steel. EPMA was used to examine this duplex scale, and showed the outermost layer to be rich in iron, chromium and oxygen, whereas only chromium and oxygen was found in the inner layer. For coated A1006 steel, no external scale was visible metallographically. A small amount of oxide that is rich in aluminium and oxygen was found by EPMA to grow inside the coating.
5.3 DISCUSSION
5.3.1 Oxidation Kinetics for Materials Tested Using TGA
At 800°C, the oxidation kinetics for A1006 steel and 2.25Cr1Mo were similar.
Their oxidation behaviours are plotted according to the parabolic rate equation and shown in Figure 5.11. It is clear the kinetics for both steels can be regarded only approximately as parabolic. Once their scales cracked or separated from the substrate, their oxidation rates were altered from the ideal diffusion controlled situation. The
2.25Cr1Mo steel showed better oxidation resistance than A06, and this may be attributed to the small addition of chromium. Chromium can promote the formation of
FeCr2O4, which has a similar structure to Fe3O4 and lower diffusion coefficient for iron.
Therefore, 2.25Cr1Mo steel has more chance to form a relatively protective Fe3O4 layer than does A1006 steel.
The unusual oxidation kinetics observed for SG cast iron can be ascribed to the combined effects of weight loss due to the consumption of graphite exposed to air and the accumulation of iron oxides. Initially, fast consumption of graphite at the surface of
SG cast iron and a relative small weight gain caused by the oxidation of iron resulted in
153 High Temperature Oxidation
a net weight loss in the initial period (Figure 5.12). When the weight gain of iron oxide exceeded the consumption of graphite, the total weight started to increase. Cavities or gaps were left after graphite was consumed or partially consumed (Figure 5.2c), but iron oxide could fill these holes and thereby provide anchors that could key the scale to the metal. Therefore, the iron oxide layer formed on SG cast iron is more protective than those formed on A06 and 2.25Cr1Mo.
Rapid weight gain kinetics were observed for the 3Cr12 steel in the initial stage and then the oxidation rate tended to slow down (second stage) (Figure 5.1b). The kinetics curve shows frequent abrupt decreases in weight during the second stage, due to scale cracking and partial spallation. The weight change kinetics (average parabolic rate is:
2x10-13 g2 cm-4 sec-1.
The stainless steels SS304, 253MA and coated A1006 steel found to have the best oxidation resistances in TGA experiments. Frequent oscillations in their weight gains
(Figure 5.1c, d) presumably reflected mechanical scale damage it is difficult to determine their oxidation kinetics after such a short period.
In long time isothermal oxidation experiments, oxidation kinetics for SS304 and coated A1006 steel were found to be parabolic (Table 5.2).
5.3.2 Scale morphology study
Duplex scale can be seen on bare A1006 steel, 2.25Cr1Mo and SG cast iron (Figure
5.2). At 800°C, all three iron oxides (Hematite, magnetite and wustite) are all stable.
According to the Fe-O binary phase diagram [207], the sequence of these iron oxides layer from the scale surface is the following: hematite/magnetite/wustite. In the case of
A1006 steel, the outmost thin layer is hematite and followed by the inner layer of
154 High Temperature Oxidation
7.0
)
-4 6.0
cm 2 5.0 (g 2 4.0 3.0 W/A) 2.0 x ( 3 1.0 10 0.0 0.0 1.0 2.0 3.0 4.0 5.0 6.0
-5 10 x Time (Sec)
(a)
1.2 ) -4 1 cm 2 0.8 (g
2 0.6
W/A) 0.4 ( x
3 0.2 10 0
0 0.5 1 1.5 2
10-5 x Time (Sec)
(b)
Figure 5.11 Diagram of weigh gain^2 vs. time for (a) A1006 steel and (b)
2.25Cr1Mo steel.
155 High Temperature Oxidation
0.1
0.05
0 )
-1 0 0.01 0.02 0.03 0.04 0.05 0.06 0.07 0.08 0.09 -0.05
-0.1
-0.15 W/A (mg cm D -0.2
-0.25
-0.3 Time (hours)
Figure 5.12 Oxidation kinetics for SG cast iron in the initial stage.
wustite. For 2.25Cr1Mo, the internal precipitate zone can also been found underneath the external scale due to the small addition of alloying elements. In Figure 5.2c, scale is found to grow underneath the consumed graphite. In cyclic oxidation, duplex scale is also formed on A1006 steel. However, the outmost hematite layer is much thicker than that formed in isothermal oxidation after a shorter exposure time.
Small amount of scale formed on chromium containing steels, such as 3Cr12,
SS304 and 253MA in both TGA and long time oxidation experiments. This may be caused by partial scale spallation and the formation of the protective Cr2O3. From their
SEM images shown in Figure 5.3, the network-like corrosion products are believed to
156 High Temperature Oxidation
be Cr2O3. In long time oxidation experiments, similar surface morphologies were observed on SS304 and 253MA. In cyclic oxidation experiments, a two-layered scale formed on SS304 steel. The scale consists of an outmost layer (FeCr2O4 spinel) and inner layer (Cr2O3). Unlike its isothermal oxidation behaviour, SS304 showed weight loss in all the cyclic oxidation experiments because of scale spallation during cooling period. The growth of the Cr2O3 layer is controlled by the outwards diffusion of Cr and inwards diffusion of O2, and the former is the more important step. As the result, cavities will be created at the interface of metal and the oxide (Figure 5.10). Therefore,
Cr2O3 layer can easily separate from the substrate. Small additions of rare earth elements, e.g. Cerium, to chromium-containing steel are known to have benefits in helping to establish the Cr2O3 quickly and slowing down its growth rate. This can explain the better oxidation resistance of 253MA. From the cross-sectional view o
253MA in Figure 5.10, only single layer scale can be seen. However, the XRD results showed that FeCr2O4 also exist. This can be attributed to the partial scale spallation.
Coated A1006 steel showed superior oxidation resistance in all the condition tested.
XRD results showed the metastable phase g, q, b and stable a phase Al2O3 formed on coated A1006 steel. It is known that the stable a-Al2O3 was normally observed around
1000°C [159]. However, the ridged structure shown in Figure 5.6 and Figure 5.7 clearly indicated that the stable a-Al2O3 formed after being exposed in air for a long time. Those needle like and plate-like oxide shown in Figure 5.6 are believed to be q-
Al2O3 [204-206]. In cyclic oxidation experiments, no evidence of the formation of iron oxide although the weight gain is slight larger than it is in isothermal oxidation experiments.
157 High Temperature Oxidation
5.4 CONCLUSION
i. Materials were tested in air at 800°C. TGA results showed that the oxidation
kinetics for all materials except SG cast iron are parabolic. For SG cast iron, the
oxidation kinetics are the combination of the consumption of graphite and
formation of iron oxide.
ii. Chromium containing steels (3Cr12, SS304 and 253MA) showed very slow
oxidation rate due to the formation of the protective Cr2O3 although partial scale
spallation occurred in TGA experiments. In the long time oxidation experiments,
partial scale spallation was observed on 253MA. In cyclic oxidation experiments,
SS304 steel showed weight loss due to the bulk scale spallation. iii. Coated A1006 steel was proved to have superior oxidation resistance in all
experiments due the formation of slow growing a-Al2O3. Aluminium containing
coating was also able to re-heal the broken scale. In long time oxidation
experiments, coated A1006 steel was found to have the similar parabolic rate to that
of SS304. In cyclic oxidation experiments, the oxidation rate of coated A1006 steel
is close to that of 253MA.
158 High Temperature Sulfidation
CHAPTER SIX
High Temperature Sulfidation
6.1 INTRODUCTION
The sulfidation resistance of the same materials tested in high temperature oxidation experiments (Chapter 5) were evaluated in SO2-containing gas mixtures at
800°C. Detailed experimental procedures were described in section 3.4.2.
6.2 THERMODYNAMIC ASPECT
As shown in Table 3.3, the gas mixture used in high temperature sulfidation experiments also contained significant amount of CO and CO2. To predict the corrosion products, the carbon activity ac was also calculated. The input gas mixture
-3 (24.9%CO+30.2%CO2+ 4.98x10 atm SO2+N2) was designed to produce a similar condition to that above an aluminium smelting cell. Calculated activities for the different species are listed in Table 6.1.
As seen in Table 6.1, the carbon activity is too low for metal carburization to take place under this condition. The sulfur containing gas species are considered to be the main reactants in addition to CO2. Therefore, the stability diagrams for metals involved in the high temperature reactions were calculated and are shown in Figure 6.1. The
-5 partial pressures for sulfur and oxygen used in the present experiments ( p = 6.54x10 S2
-19 atm, p =3.02x 10 atm) are marked as point A in the diagram. From the stability O2 diagram, it is found that aluminium and chromium are stable as alumina and chromia
159 High Temperature Sulfidation
Table 6.1 Equilibrium partial pressure and activities for species in high temperature
sulfidation experiments at 800°C.
Species Partial Pressure (atm)/Activity
O2 (g) 3E-19
CO2 (g) 2.6E-01
CO (g) 2.9E-01
C (g) 8.6E-29
COS (g) 4.8E-03
CS2 (g) 2.1E-04
S2 (g) 6.4E-05
SO2 (g) 1.7E-07
SO3 (g) 6.7E-17
N2 (g) 4.5E-01
C (s) 4.6E-02
160 High Temperature Sulfidation Stability diagram of M-S-O system at 800°C. at system Figure 6.1 Stability diagram of M-S-O
161 High Temperature Sulfidation
while iron and nickel are stable as iron sulfide and nickel sulfide. However, the oxygen potential is high enough to form FeO and perhaps Fe3O4 in the absence of sulfur.
6.3 RESULTS
Materials were exposed to the SO2-containing gas mixture for different periods in order to determine their sulfidation kinetics. Weight gains of materials are plotted against the exposure time and shown in Figure 6.2.
According to their weight gains, the sulfidation resistance of materials is characterized into two groups. Bare A1006 steel, 2.25Cr1Mo, SG cast iron, 3Cr12 and
SS304 were found to have the worst sulfidation resistance among all the materials.
However, only relatively small amounts of corrosion products were measured on
253MA and coated A1006 steel. As expected, A1006 steel showed the worst sulfidation resistance, followed by 2.25Cr1Mo. SG cast iron showed somewhat better sulfidation resistance and this may have been caused by the blockage effect of the graphite. The weight gains of SS304 were found to be slightly higher than those of
3Cr12. The sulfidation kinetics for these materials were plotted according to the parabolic rate equation and are shown in Figure 6.3. The calculated parabolic rates are listed in Table 6.2.
Thick, grey coloured scale could be seen on bare A1006 steel, 2.25Cr1Mo, SG cast iron, 3Cr12, and SS304 stainless steel. A relatively thin, grey coloured scale was formed on 253MA steel. No significant amount of scale was formed on coated A1006 steel.
The corrosion products formed on these materials were then analysed using XRD, and the results are shown in Table 6.3. As expected, iron sulfides and magnetite were
162 High Temperature Sulfidation
120 2.25Cr1Mo A06 100 SG cast iron SS304 80 ) 3Cr12 2
60 W/A (mg/cm
∆ 40
20
0 0481216 Time (hours)
2.5 253MA Coated A06 2 ) 2 1.5
1 W/A (mg/cm ∆
0.5
0 0481216 Time (hours)
Figure 6.2 Weight gain of samples in sulfidation experiments at 800°C.
163 High Temperature Sulfidation
1.4 2.25Cr1Mo 1.2 ) A06 -4 SG cast iron
cm 1 2 0.8 x mg 4
(10 0.6 2 0.4 W/A) ∆ ( 0.2
0 01234567 10-4 Time (Sec) Figure 6.3a
7
6 3Cr12 )
-4 SS304 5 cm 2 4 x mg 3 3 (10 2 2 W/A) ∆
( 1
0 01234567
10-4 Time (Sec)
Figure 6.3b
164 High Temperature Sulfidation
5 253MA 4 Coated A06 ) -4 3 cm 2
(mg 2 2 W/A)
∆ 1 (
0 01234567
10-4 Time (Sec)
Figure 6.3c
Figure 6.3 Kinetics of materials in sulfidation experiments at 800°C.
Table 6.2 Parabolic sulfidation rates at 800°C.
2 -4 -1 2 -4 -1 Materials kp (g .cm .sec ) Materials kp (g .cm .sec )
A06 2.13 x 10-7 SS304 9.80 x 10-8
2.25Cr1Mo 1.82 x 10-7 253MA 7.57 x 10-11
SG cast iron 0.95 x 10-7 Coated A06 0.21 x 10-11
3Cr12 6.72 x 10-8
165 High Temperature Sulfidation
Table 6.3 XRD results of in situ scales grown in high temperature
sulfidation experiments.
Samples XRD Results
A06 Fe1-xS, Magnetite
2.25Cr1Mo Fe1-xS, Magnetite, FeCr2O4
SG cast iron Fe1-xS, Magnetite
3Cr12 Fe1-xS, Magnetite, FeCr2O4
SS304 Fe1-xS, Magnetite, FeCr2O4
253MA Fe1-xS, FeCr2O4, Cr2O3
Coated A06 α−Al2O3
166 High Temperature Sulfidation
found on A1006 steel and SG cast iron. Iron chromium spinel was found on
2.25Cr1Mo, 3Cr12 and SS304 steel. However, no Cr2O3 could be found on these three chromium containing alloys. This may explain their worse sulfidation resistance under this condition. The better sulfidation resistance of 253MA and coated A1006 steel can be attributed to the formation of chromia and α−Al2O3, respectively. No iron sulfide was found on coated A1006 steel.
Reacted samples were then cross-sectioned after different exposure periods. Their cross-sectional optical images are shown in Figure 6.4 - 6.10. EDAX was also employed to analyse the scale constitution.
The scale formed on A1006 steel sulfidised after 2 hr was composed of a thin dark coloured outermost layer, an intermediate layer containing both FeS (light phase) and iron oxide (grey phase) and an thin grey coloured innermost layer. Gaps were found at the interface of outermost layer/intermediate and scale/metal. With increasing exposure time, the scale structure formed after 4 hr consisted of a thin, grey coloured outmost oxide layer and an intermediate layer with the duplex structure and a grey coloured innermost layer with light FeS dispersion. After 6 hr, the scale can be regarded as a single duplex layer. For A1006 steel sulfidised after 8 and 16hr, the duplex structure was found to constitute the outermost layer, under which was a thin, grey coloured innermost oxide layer.
Scale formed on 2.25Cr1Mo steel after 2 hr exposure consisted of a top grey coloured iron oxide layer followed by a thick layer of iron sulfide plus iron oxide duplex structure and an innermost layer of metal sulfide(s) in the iron oxide matrix.
The thin outermost Fe3O4 layer was not found for 2.25Cr1Mo steel after 4 hr and 6 hr reaction. Instead, wustite was the outermost layer followed by an internal zone of iron
167 High Temperature Sulfidation
oxide layer containing precipitate of metal sulfide. This layer was also found to be rich in chromium and silicon, using EDAX. The lamellar duplex structure was observed again for samples after 8 hr. The scale formed consisted of a thin oxide layer followed by a thick lamellar structure scale, and an innermost fine-grained, multiphase zone.
After 16 hours, the outermost oxide layer mentioned above was not found. The scale structure consisted of an outermost duplex structure and an innermost fine-grained multiphase layer.
The duplex structure was not observed on SG cast iron after the initial 2 hours exposure (Figure 6.6). Instead, the scale consisted of a thin outermost Fe3O4 layer, an intermediate layer of wustite and an internal fine-grained, multiphase zone. Un-reacted graphite was found innermost zone. After 4 hours, the outermost layer formed on SG cast iron was FeO only. The incorporation of graphite into the scale is clearly seen in the cross-section of a sample reacted for 8h. Scales formed on SG cast iron after 6, 8,
16 hours were similar: a thin magnetite layer followed by a thick layer of lamellar structure and an innermost multiphase zone. EDAX analysis showed the innermost layer to be enriched in sulfur, oxygen and silicon.
Compared to the scales formed on 2.25Cr1Mo and SG cast iron, a similar but relatively thin scale formed on 3Cr12 (Figure 6.7). After 2hr exposure, a thin, grey coloured iron oxide layer was found to be the outermost layer, followed by a porous iron sulfide layer and an innermost fine-grained, multiphase layer. The topmost thin iron oxide layer could not be seen in some cases and is attributed to partial scale spallation. The innermost zone was found to be rich in chromium, sulfur and oxygen besides iron. Large numbers of cracks were found to cross the whole scale thickness, and this may explain the poor sulfidation resistance of 3Cr12.
168 High Temperature Sulfidation
Figure 6.4a
Figure 6.4b
169 High Temperature Sulfidation
Figure 6.4c
Figure 6.4d
170 High Temperature Sulfidation
Figure 6.4e
Figure 6.4 Cross-sectional images for A1006 steel in high temperature sulfidation
experiments at 800°C after different exposure time.
171 High Temperature Sulfidation
Figure 6.5a
Figure 6.5b
172 High Temperature Sulfidation
Figure 6.5c
Figure 6.5d
173 High Temperature Sulfidation
Figure 6.5e
Figure 6.5 Cross-sectional images for 2.25Cr1Mo steel in high temperature
sulfidation experiments at 800°C after different exposure time
Metallography studies for SS304 showed two-layered structures: an external iron sulfide layer and an inner layer consisting of fine oxide particles inside a sulfide matrix
(Figure 6.8). EDAX analysis showed the innermost layer also contains small amount of nickel and trace of chromium. It was also found that the sample corner was attacked more severely than other sample areas.
For 253MA after 2 hours exposure, only a small amount of grey coloured oxide scale was found, due to partial scale spallation (Figure 6.9). After reaction of 253MA steel for up to 8 hours, a two-layered scale consisted of an outermost layer of iron oxide and iron sulfide and an inner mixed zone. After 16 hours, a thin oxide layer followed
174 High Temperature Sulfidation
by a relative thick iron sulfide layer formed the external scale region. EDAX analysis of its innermost zone showed it to be similar to those of 3Cr12 or SS304: consisting of nickel and chromium, oxygen and sulfur besides iron.
As seen in Figure 6.10, no obvious corrosion products could be found in the cross- sectional images of coated A1006 steel after different exposure time. However, EDAX analysis on the coating surface revealed the scale on coated A06 was rich in aluminium and oxygen. No sulfur was found at the surface.
A surface morphology study was carried out using FESEM on materials after 16 hours exposure. The appearance of the scale outer surfaces is shown in Figure 6.11.
The surface morphologies found for A1006, 2.25Cr1Mo and SG cast iron were similar.
The top porous scale surface was found to be rich in iron and oxygen. Some iron oxide tips were also found. Scale formed on 3Cr12 was similar to those formed on the previous three materials, reflecting the formation of the same corrosion product (iron oxide), but without any oxide whiskers or outgrowths. The scale formed on SS304 was found to be similar to that of 3Cr12, no protective scale can be seen as well. In the case of 253MA steel, the outer scale was relatively compact. In Figure 6.12, corrosion product protrusions are seen to have grown through the scale surface. EDAX results showed that these corrosion products were rich in nickel and sulfur and they are believed to be nickel sulfide. EDAX analysis revealed the scale on coated A06 to be enriched in aluminium and oxygen. The surface morphology of the sulfidation reaction product was very similar to that developed on coated A06 in oxidation experiments
(Figure 5.7a). Large numbers of small ridges formed were believed to be α-Al2O3. The
EDAX result for the coated A06 is in good agreement with the XRD result which showed only α-Al2O3 formed.
175 High Temperature Sulfidation
Figure 6.6a
Figure 6.6b
176 High Temperature Sulfidation
Figure 6.6c
Figure 6.6d
177 High Temperature Sulfidation
Figure 6.6e
Figure 6.6 Cross-sectional images for SG cast iron in high temperature sulfidation
experiments at 800°C after different exposure time
Figure 6.7a
178 High Temperature Sulfidation
Figure 6.7b
Figure 6.7c
179 High Temperature Sulfidation
Figure 6.7d
Figure 6.7e
Figure 6.7 Cross-sectional images for 3Cr12 steel in high temperature sulfidation
experiments at 800°C after different exposure time
180 High Temperature Sulfidation
Figure 6.8a
Figure 6.8b
181 High Temperature Sulfidation
Figure 6.8c
Figure 6.8d
182 High Temperature Sulfidation
Figure 6.8e
Figure 6.8 Cross-sectional images for SS304 steel in high temperature sulfidation
experiments at 800°C after different exposure time
183 High Temperature Sulfidation
Figure 6.9a
Figure 6.9b
184 High Temperature Sulfidation
Figure 6.9c
Figure 6.9d
185 High Temperature Sulfidation
Figure 6.9e
Figure 6.9 Cross-sectional images for 253MA steel in high temperature sulfidation
experiments at 800°C after different exposure time.
186 High Temperature Sulfidation
Figure 6.10a
Figure 6.10b
187 High Temperature Sulfidation
Figure 6.10c
Figure 6.10d
188 High Temperature Sulfidation
Figure 6.10e
Figure 6.10 Cross-sectional images for coated A1006 in high temperature sulfidation
experiments at 800°C after different exposure time.
189 High Temperature Sulfidation
A06 2.25Cr1Mo
3 µm 3 µm
SG cast iron 3Cr12
µ 3 m 6 µm
SS304 253MA
3 µm 6 µm
Coated A06 Coated A06
3 µm 20 µm
Figure 6.11 SEM images for materials sulfidised at 800°C for 16 hours.
190 High Temperature Sulfidation
Figure 6.12 SEM image for 253MA sulfidised at 800°C for 16 hours.
6.4 DISCUSSION
For A1006 steel, 2.25Cr1Mo and SG cast iron, duplex structures of iron sulfide and iron oxide were found. This is in conflict with the thermodynamic prediction of forming iron sulfide phase only (Figure 6.1). The same observation has been repeated many times [3, 208-212] for reaction of pure iron with SO2 gases. The scale micro structures for these materials evolved differently with time.
In the case of A1006 steel, the scale morphologies for samples after different exposure times were all found to have the duplex structure. However, the thin innermost and the outermost iron oxide layer were sometimes not found in metallographic cross-sectional images. Obviously, the equilibrium between the input bulk gas phase and the scale surface was not achieved. Therefore, SO2 itself is the
191 High Temperature Sulfidation
likely reactant. This issue has been widely investigated and repeatedly reported by many researchers [3, 208-212]. The following reactions have been proposed to explain the way in which of SO2 interacted with the scale surface:
3Fe + SO2 = FeS + 2FeO (6.1)
5 1 Fe + SO2 = FeS + Fe3O4 (6.2) 2 2
Fe3O4 + 3SO2 = 3FeS + 5O2 (6.3)
5 3FeS + 2SO2 = Fe3O4 + S2 (6.4) 2
1 2Fe + SO2 = 2FeO + S2 (6.5) 2
XRD results for A1006 steel revealed that the iron oxide was Fe3O4. Therefore, equation 6.2 is believed to be responsible for the formation of the duplex scale structure.
Reactions (6.3) – (6.5) are either thermodynamically not favoured, or too slow to contribute
Of course, this kind of structure could not provide protection for these steels. The innermost oxide layer with fine dispersed iron sulfide implied a significant inward transport of sulfur. The iron sulfide dispersion was found at the outer portion of the innermost scale but not at the metal/scale interface because sulfur activity was too low to form iron sulfide even though the metal was at unit activity. When SO2 is the reactant, it is easy to explain the formation of the occasionally observed (after 4hr exposure) outermost iron oxide layer (mainly Fe3O4) according to equations 6.2 and 6.4.
The sulfidation rate for A1006 steel obtained in this study was found to be similar to that of pure iron which was found to be 1.3E-7 g2.cm-4.sec-1 in the gas mixtures containing SO2-CO-CO2-N2 at the same temperature [212].
192 High Temperature Sulfidation
The 2.25Cr1Mo steel behaved similarly to the low carbon steel, except that it developed an innermost layer which was enriched in chromium and silicon. Although the evidence was incomplete, it is concluded that the stable phases FeCr2O4 and SiO 2
(or perhaps FeSiO 3) formed in the inner layer as discrete particles. The heterophase inner layers would be expected to possess low plasticity. This kind of scale could not provide a good barrier between metal and reaction gas. Therefore, it was likely that the reaction gas could penetrate through those defects to reach the metal surface. This would explain the formation of sulfides particles in the innermost layer.
Similar to 2.25Cr1Mo exposed for 4 and 6 hours, SG cast iron formed a thin magnetite layer followed by a thick FeO intermediate layer and an inner multi-phase zone after 2 hours. The scale was found to be not protective, and cracks and gaps developed through the thickness of the outermost layer. Oxidant could easily penetrate through the scale. After 4 hours, the thin outermost magnetite layer observed previously was not found, probably as a result of the partial scale spallation. Duplex oxide plus sulfide structures also grew on SG cast iron, showing that the equilibrium between bulk gas phase and scale surface was not reached. However, SG cast iron showed somewhat better corrosion resistance, and this may have been caused by the blockage effect of the graphite.
According to the stability diagram for M-S-O systems (Figure 6.1), chromium and aluminium are stable in the form of metal oxide in the gas mixture used in these high temperature sulfidation experiments. Therefore, the chromium-containing steels were supposed to be protected by chromia. However, no chromia could be found (using
XRD) on these two samples after different exposure times. The formation of a thin topmost oxide layer and an underlying iron sulfide layer indicate that equilibrium
193 High Temperature Sulfidation
between bulk gas phase and scale surface was not reached. Similar parabolic rates were observed for 3Cr12 and SS304 steel, indicating that the higher chromium level did not supply a better sulfidation resistance for SS304 steel. In the case of the oxidation of chromium containing steel, the chromium level determines the oxidation resistance of chromium containing steel. As seen in Fe-Cr-O phase diagram (Figure 6.13), chromium
* containing steel cannot form a permanent protective chromia scale unless N Cr is
2+ exceeded [8]. Increasing the chromium level, Fe ion will be blocked by the FeCr2O4 islands and the thickness of FeO layer is thinned. However, the diffusibility of iron ion inside chromia scale is still fast enough to form an outer iron oxide dominated layer even at the higher chromium level. In the sulfidation conditions, the circumstance is more complex. It is known that the sulfidation resistance of steel containing intermediate chromium level (12wt%) is twice as good as carbon steel in H2/H2S atmosphere at around 900K [90-92 95]. Although different gas mixtures and temperature was used in the present study, the corrosion rates of these two materials are found to be twice or three times as good as low carbon A1006 steel. These two materials are also very similar in their scale constitution and morphology. In their cross-section view, an outermost oxide layer can be seen sometimes on 3Cr12 but not on SS304. However, the SEM analysis for their scale surface showed same structure.
Compared to 2.25Cr1Mo, the better sulfidation resistance of these two materials may be caused by the progressive replacement of iron with chromium in the innermost spinel leading to a lower diffusivity iron in this layer. Hence, the growth of the outermost layer was limited. Cracks, pores and gaps between and at the interface of the outermost layer/innermost layers became the dominant scale defects on these two samples.
Therefore, there was more ready access for oxidant through these scale defects to attack
194 High Temperature Sulfidation
the metal underneath the porous outermost scale. This may explain their bad sulfidation resistances.
Compared to 3Cr12 and SS304 steel, 253MA heat resistant steel showed much better sulfidation-oxidation resistance. The thin oxide layer observed for 253MA after
16 hours exposure was found to be rich in chromium and oxygen using EDAX. XRD results confirmed the formation of chromia. Partial scale spallation made analysis of this scale difficult. Although a small amount of iron sulfide was found by XRD on
253MA steel, no obvious iron sulfide is seen in the SEM images of the external surface
S1=Fe1.5Cr1.5O4
S2=FeCr2O4
O (at%) Fe1-xO+S1
Fe2O3 Cr2O3 Fe3O4 Fe O 1-x O 3 1 1 2 + S O+S γ + Cr x 2 1- + S γ γ -Fe + Fe1-xO -Fe + Fe γ α -Fe + Cr2O3
3 O α γ 2 + +Cr2O3 γ+ Cr
Cr (at%)
Figure 6.13 Isothermal section for the Fe-Cr-O phase diagram at 1200°C [8].
195 High Temperature Sulfidation
shown in Figures 6.11 and Figure 6.12. It is known that the formation of iron sulfide is mainly controlled by the outwards diffusion of iron [213]. However, the rate at which iron sulfide grows on pure iron is much faster than the slow rates observed on 253MA.
Therefore, it is believed that the growth rate of iron sulfide was controlled by the relatively slow diffusion of iron through a reasonably compact chromia layer. The formation of the chromia layer on 253MA may be attributed to its high chromium level and also the addition of the rare earth element (Cerium). However, there is no direct evidence showing the beneficial effect of Cerium due to scale spallation of 253MA during cooling. At 1073K, nickel sulfide is stable as a liquid [214]. The liquid Ni-S phase could penetrate the corrosion product layer and form protrusion of the corrosion product (Figure 6.12) with a NiO shell and nickel sulfide core [215]. This potentially damaging process was restricted in development by the existence of a relatively protective chromia layer.
It is known that Al2O3 scales are much better barriers to sulfur-containing environments than Cr2O3 scales [216], by virtue of their superior diffusion properties.
A study of the sulfidation resistance of iron aluminides (containing 16-40at% aluminium) at 800°C showed that the threshold aluminium level for optimum sulfidation resistance was 18at% aluminium [178]. It was also found that the scales formed on those alloys which had better sulfidation resistance were composed totally of
Al2O3. In the current case, the aluminium level in the coating was more than the threshold level. Phase identification by XRD analysis and surface morphological studies confirmed that a slow growing and compact α-Al2O3 layer formed on coated
A1006 steel. This α-Al2O3 layer then acted as a barrier between base metal and gas mixture. Therefore, it is not surprised that superior sulfidation resistance was achieved.
196 High Temperature Sulfidation
6.5 CONCLUSION
i. Different materials were tested in a low oxygen activity, sulfur-containing gas
mixture at 800°C. Fast scale formation was observed on A1006, 2.25Cr1Mo,
steel and SG cast iron. Their rapid corrosion was caused mainly by the
formation of fast growing iron sulfide phase. SG cast iron showed somewhat
better sulfidation resistance, which can be attributed to the blockage effect of
graphite incorporated into the scale from SG cast iron. Scale formation on these
materials was controlled by both outwards diffusion of iron and inwards
diffusion of oxidants. Reaction kinetics were consequently parabolic.
ii. Unlike 253MA heat resistant steel, chromium-containing steels (3Cr12 and
SS304) showed unacceptable sulfidation resistance due to the formation of fast
growing iron sulfide. The formation of a protective chromia layer was observed
only on 253MA. Scale formation was also controlled by both outwards
diffusions of metal and inwards diffusion of oxidants. Parabolic kinetics were
found for all three materials, the slower reaction rate of 253MA reflecting the
slow diffusion iron through its protective Cr2O3 scale layer.
iii. Coated A1006 steel was shown to have superior oxidation resistance in all
experiments, due to the formation of slow growing α-Al2O3. The alumina
scales formed on coated A1006 steel were rather compact and therefore blocked
the contact between base metal and gas mixture. Alumina formation was as
predicted for the high aluminium content coating.
iv. The equilibrium at the interface of bulk gas mixture and metal surface was not
achieved although catalyst was used. Therefore, the classic duplex structure
was found in most cases.
197 Low Temperature Corrosion Tests in Gas Mixtures
CHAPTER SEVEN Low Temperature Corrosion Tests
7.1 INTRODUCTION
Materials used in aluminium smelting cell structures are also required to resist the low temperature corrosion by gases containing SO2 and HF. To determine the effects on their corrosion resistances, materials were tested under four different gas mixtures.
Instead of using HF, HCl gas was substituted out of regard for laboratory safety. The gas mixtures 1-4 are defined in Table 3.3.
7.2 THERMODYNAMIC ASPECTS
In addition to hydrogen chloride gas, sulfur dioxide was also used in low temperature corrosion tests. Partial pressures for the various possible gas species were calculated using Chemix and are listed in Table 7.1. Stability diagrams for metal- chlorine-oxygen (M-Cl-O) and metal-sulfur-oxygen (M-S-O) were also calculated to predict the corrosion products (Figure 7.1). The experimental gas compositions are also marked in these diagrams. As seen in Figure 7.1a,and 7.1b, all metal oxides are stable with respect to chlorides when in contact with the reaction gases, although the gas chlorine potential was above the equilibrium values for metal chloride production. In the gas containing sulfur dioxide, most metals are predicted to form sulphate, due to the extremely low partial pressure of sulphur, but significant pressure of SO3. However, these predictions are based of course on an assumed equilibrium at the interface of the bulk gas and solid surface. To equilibrate the gases, they were passed over alumina- supported platinum catalyst in all the experiments.
198 Low Temperature Corrosion Tests in Gas Mixtures
Table 7.1 Equilibrium partial pressure and activities for species in high temperature
sulfidation experiments at 400°C
Gas 1 Gas 2 Gas 3 Gas 4
(Air) (Air+H2O) (Air+HCl+H2O) (Air+HCl+H2O+SO2)
Species Partial Pressure Partial Pressure Partial Pressure Partial Pressure
(atm) (atm) (atm) (atm)
O2 (g) 2.1E-01 2.0E-01 2.0E-01 2.1E-01
N2 (g) 7.8E-01 7.6E-01 7.5E-01 7.5E-01
HCl (g) - 4.4E-04 4.5E-04
H2O (g) 3.7E-02 3.7E-02 3.8E-02
H2 (g) 4.0E-18 4.0E-18 4.0E-18
Cl2 (g) - 2.5E-05 2.6E-05
Cl (g) - 1.4E-09 1.5E-09
HClO (g) - 3.0E-07 3.0E-07
S2 (g) - - 7.3E-59
SO2 (g) - - 3.3E-06
SO3 (g) - - 8.5E-04
199 Low Temperature Corrosion Tests in Gas Mixtures
Figure 7.1a
Figure 7.1b
200 Low Temperature Corrosion Tests in Gas Mixtures
Figure 7.1C
201 Low Temperature Corrosion Tests in Gas Mixtures
Figure 7.1d
Figure 7.1 Stability diagrams for materials tested in M-Cl-O and M-S-O systems
at 400°C .
202 Low Temperature Corrosion Tests in Gas Mixtures
7.3 RESULTS
All materials were exposed at 400°C for different periods. Their weight gains were measured after exposure and are shown in Figure 7.2.
As seen in Figure 7.2, the corrosion performance of the different materials varied widely. The weight changes for corrosion-resistant materials (253MA, SS304, aluminised steel and aluminium and zincalume) very irregular. The irregular weight changes may be attributed to partial scale spallation. 3Cr12 was found to react more rapidly in the SO2-containing gas mixture than did SS304 and 253MA, although its weight changes were still very small. Pure copper and Cu-S alloys were found to have the worst corrosion resistance among all the materials tested, especially in SO2- containing gas mixture. Bare A1006 steel, 2.25Cr1Mo and SG cast iron formed a group exhibiting similar weight changes. This group of materials all showed worse corrosion resistance in SO2-containing gas.
To demonstrate the effects of different gas compositions on materials, weight changes for samples exposed at 400°C for 500 hours are shown in Figure 7.3. The occasions where no weight change could be measured are specified in the diagram.
The appearance of the reacted samples varied considerably. The scales formed on
A1006 steel, 2.25Cr1Mo, and SG cast iron in all of the different gas mixtures were similar: thin and black coloured. No obvious scale could be found on 3Cr12, SS304,
253MA, pure aluminium, zincalume, and coated A1006 steel after reaction with any of the gases. For pure copper and Cu-S alloy, thin, black coloured scales were formed in air and air with water vapour. Extensive scale spallation occurred on these two materials during cooling. Thick, grey coloured scales were observed when HCl was added to the gas.
203 Low Temperature Corrosion Tests in Gas Mixtures
A06 steel 3.5 Gas 1 3 Gas 2 Gas 3 2.5 Gas 4 ) 2 2
1.5 W/A (mg/cm 1
0.5
0 0 100 200 300 400 500 600 Time (hours)
Figure 7.2a
2.25Cr1Mo steel 4.5 Gas 1 4 Gas 2 3.5 Gas 3 ) 2 3 Gas 4 2.5 2
W/A (mg/cm 1.5 1 0.5 0 0 100 200 300 400 500 600
Time (hours)
Figure 7.2b
204 Low Temperature Corrosion Tests in Gas Mixtures
SG cast iron
5 Gas 1 4 Gas 2 Gas 3
) Gas 4 2 3
2 W/A (mg/cm 1
0 0 100 200 300 400 500 600 -1
Time (hours)
Figure 7.2c
3Cr12
1 Gas 1 0.8 Gas 2 Gas 3 0.6 Gas 4 ) 2 0.4
0.2 W/A (mg/cm
0 0 100 200 300 400 500 600 -0.2
-0.4 Time (hours)
Figure 7.2d
205 Low Temperature Corrosion Tests in Gas Mixtures
SS304 0.4 Gas 1 0.3 Gas 2 Gas 3 0.2 )
2 Gas 4 0.1
0 W/A (mg/cm 0 100 200 300 400 500 600 -0.1
-0.2
-0.3
Time (hours)
Figure 7.2e
253MA 0.4 Gas 1 0.3 Gas 2 Gas 3 ) 2 0.2 Gas 4
0.1 W/A (mg/cm
0 0 100 200 300 400 500 600 -0.1
-0.2
Time (hours)
Figure 7.2f
206 Low Temperature Corrosion Tests in Gas Mixtures
Coated A06 0.3 0.25 0.2 0.15 ) 2 0.1 0.05 0
W/A (mg/cm -0.05 0 100 200 300 400 500 600 -0.1 Gas 1 Gas 2 -0.15 Gas 3 -0.2 Gas 4 -0.25
Time (hours)
Figure 7.2g
Pure Al 0.25 Gas 1 0.2 Gas 2 0.15 Gas 3 Gas 4 0.1 ) 2 0.05 0 0 100 200 300 400 500 600
W/A (mg/cm -0.05 -0.1 -0.15 -0.2
-0.25 Time (hours)
Figure 7.2h
207 Low Temperature Corrosion Tests in Gas Mixtures
Zincalume 0.2
0.1
0 ) 2 0 100 200 300 400 500 600 -0.1
-0.2 W/A (mg/cm Gas 1 -0.3 Gas 2 Gas 3 -0.4 Gas 4
-0.5 Time (hours) Figure 7.2i
Pure Cu 25 Gas 1 Gas 2 20 Gas 3 ) 2 Gas 4 15
10 W/A (mg/cm
5
0 0 100 200 300 400 500 600
Time (hours) Figure 7.2j
208 Low Temperature Corrosion Tests in Gas Mixtures
Cu-S alloy 80 Gas 1 70 Gas 2 60 Gas 3
) Gas 4 2 50
40
W/A (mg/cm 30
20
10
0 0 100 200 300 400 500 600 Time (hours)
Figure 7.2k
Figure 7.2 Kinetics for materials tested at 400°C in different gas mixtures.
209 Low Temperature Corrosion Tests in Gas Mixtures
The corrosion products formed on these materials were then analysed using XRD.
The results are shown in Table 7.2. For A1006, 2.25Cr1Mo and SG cast iron, only oxides were formed in the gas mixtures without SO2. The gas containing SO2 formed the same oxides and, in addition, Fe1-xS. The greatly increased weight gains accompanying exposure to Gas 4 is therefore attributed to the formation of non- protective iron sulfide product. No corrosion products could be found using XRD on
3Cr12, SS304 and 253MA exposed to gas mixtures without any SO2 content. This was due to the only negligible amount of corrosion product formed on these materials. In the case of 3Cr12, Fe1-xS was found in the gas with SO2, thereby explaining its poor corrosion resistance. The stainless steels SS304 and 253MA revealed no corrosion products after exposure to any of the gases, consistent with their excellent corrosion resistance. Metastable alumina phases were found on coated A1006 steel whereas no corrosion products were detected on pure aluminium, despite the fact that weight changes produced on the two materials were similar. Oxides were found on zincalume in all the tests. Copper oxides were formed on both copper and Cu-S alloy in reaction with air and air with water vapour. After exposure to gas containing HCl, trace amounts of CuCl were also found. A small amount of copper sulfide was also formed when SO2 was added to the gas.
Metallographic images of some materials after exposure to reaction gas for 500 hours are shown in Figure 7.4. For materials such as 3Cr12, SS304 and 253MA, no observable corrosion product could be found after reaction with any of the gas mixtures.
Cross-sectional images for these materials are therefore not shown. There was no observable corrosion products found in cross-sectional views of zincalume and aluminium after reaction in Gases 1 and 2. Therefore, only the cross-sectional images of zincalume and aluminium after exposure to Gases 3 and 4 for 500 hours are also
210 Low Temperature Corrosion Tests in Gas Mixtures
Table 7.2 XRD results of materials reacted in different gas mixtures at 400°C
Gas 1 Gas 2 Gas 3 Gas 4
Air Air+H2O Air+H2O+HCl Air+H2O+HCl+SO2
Fe1-xS Fe2O3 Fe2O3 Fe2O3 A06 Fe2O3 Fe3O4 Fe3O4 Fe3O4 Fe3O4
Fe2O3 Fe2O3 Fe1-xS Fe2O3 Fe3O4 Fe3O4 Fe2O3 2.25Cr1Mo Fe3O4 Cr-O Cr2O3 Fe3O4 FeCr2O4 FeCr2O4 FeCr2O4 FeCr2O4
Fe2O3 SG Fe2O3 Fe2O3 Fe2O3 Fe3O4 cast iron Fe3O4 Fe3O4 Fe3O4 Fe1-xS
Fe1-xS
Fe2O3
3Cr12 Ferrite Ferrite Ferrite Fe3O4
FeCr2O4
Ferrite
SS304 Austenite Austenite Austenite Austenite
253MA Austenite Austenite Austenite Austenite
ι, β-Al2O3 ι-Al2O3 ι-Al2O3 ι-Al2O3 Coated FeAl2 FeAl2 FeAl2 FeAl2 A1006 FeAl, Fe3Al FeAl, Fe3Al FeAl, Fe3Al FeAl, Fe3Al
211 Low Temperature Corrosion Tests in Gas Mixtures
Al Al Al Al Al
ZnO ZnO ZnO ZnO Zincalume γ -Al2O3 γ -Al2O3 γ -Al2O3 γ -Al2O3
CuO CuO CuO CuO Cu2O Cu Cu2O Cu2O Cu2O CuCl CuCl Cu2S
CuO CuO CuO CuO Cu2O Cu-S alloy Cu2O Cu2O Cu2O CuCl CuCl Cu2S
5 2.25Cr1Mo 4.5 Bare A06 4 SG cast iron 3.5 ) 2 3
2.5
2 W/A (mg/cm
1.5
1
0.5
0 Gas 1 Gas 2 Gas 3 Gas 4
Figure 7.3a
212 Low Temperature Corrosion Tests in Gas Mixtures
1 253MA 0.8 3Cr12 Coated A06 0.6 SS304 ) 2
0.4
W/A (mg/cm 0.2
0.00 0 0.00 Gas 1 Gas 2 Gas 3 Gas 4 -0.2
-0.4
Figure 7.3b
0.1
0.05 0.00 0 Gas 1 Gas 2 Gas 3 Gas 4 -0.05 ) 2 -0.1
-0.15
W/A (mg/cm -0.2
-0.25
-0.3 Pure Al -0.35 Znicalume
-0.4
Figure 7.3c
213 Low Temperature Corrosion Tests in Gas Mixtures
80 Cu 70 Cu-S
60 ) 2 50
40 W/A (mg/cm 30
20
10
0 Gas 1 Gas 2 Gas 3 Gas 4 Figure 7.3d
Figure 7.3 Weight changes for materials exposed to different gas mixtures for 500h at
400 °C. shown for comparison.
In Gas 1, the scales formed on A1006 steel, 2.25Cr1Mo and SG cast iron were similar except in their thickness: two layered structures consisting of a thin outermost layer of Fe2O3 above a relatively thick inner Fe3O4 layer (Figure 7.4a, 7.4e and 7.4i). A porous, two-layered structure consisted of a topmost thin, dark grey coloured CuO layer and a relatively thick inner Cu2O layer was observed on pure copper and Cu-S alloy
(Figure 7.4q and 7.4u). Within the Cu-S alloy, preferential attack had occurred at the site of the original copper sulfide inclusions. In the case of the coated A1006 steel, no observable corrosion products could be found in its cross-sectional view (Figure 7.4m)
214 Low Temperature Corrosion Tests in Gas Mixtures
Figure 7.4a A1006 exposed to gas 1 for 500 hours
Figure 7.4b A1006 exposed to gas 2 for 500 hours
Figure 7.4c A1006 exposed to gas 3 for 500 hours
215 Low Temperature Corrosion Tests in Gas Mixtures
Figure 7.4d A1006 exposed to gas 4 for 500 hours
Figure 7.4e 2.25Cr1Mo exposed to gas 1 for 500hours
Figure 7.4f 2.25Cr1Mo exposed to gas 2 for 500hours
216 Low Temperature Corrosion Tests in Gas Mixtures
Figure 7.4g 2.25Cr1Mo exposed to gas 3 for 500hours
Figure 7.4h 2.25Cr1Mo exposed to gas 4 for 500hours
Figure 7.4i SG cast iron exposed to gas 1 for 500hours
217 Low Temperature Corrosion Tests in Gas Mixtures
Figure 7.4j SG cast iron exposed to gas 2 for 500hours
Figure 7.4k SG cast iron exposed to gas 3 for 500hours
Figure 7.4l SG cast iron exposed to gas 4 for 500hours
218 Low Temperature Corrosion Tests in Gas Mixtures
Figure 7.4m Coated A1006 exposed to gas 1 for 500hours
Figure 7.4n Coated A1006 exposed to gas 2 for 500hours
Figure 7.4o Coated A1006 exposed to gas 3 for 500hours
219 Low Temperature Corrosion Tests in Gas Mixtures
Figure 7.4p Coated A1006 exposed to gas 4 for 500hours
Figure 7.4q Copper exposed to gas 1 for 500hours
Figure 7.4r Copper exposed to gas 2 for 500hours
220 Low Temperature Corrosion Tests in Gas Mixtures
Figure 7.4s Copper exposed to gas 3 for 500hours
Sulfide chloride
Figure 7.4t Copper exposed to gas 4 for 500hours
Figure 7.4u Cu-S alloy exposed to gas 1 for 500hours
221 Low Temperature Corrosion Tests in Gas Mixtures
Figure 7.4v Cu-S alloy exposed to gas 2 for 500hours
Figure 7.4w Cu-S alloy exposed to gas 3 for 500hours
Sulfide chloride
Figure 7.4x Cross-section image for Cu-S alloy exposed to gas 4 for 500hours
222 Low Temperature Corrosion Tests in Gas Mixtures
Figure 7.4y Zincalume exposed to gas 3 for 500hours
Figure 7.4z Zincalume exposed to gas 4 for 500hours
Figure 7.4aa Aluminium exposed to gas 3 for 500hours
223 Low Temperature Corrosion Tests in Gas Mixtures
Figure 7.4ab Aluminium exposed to gas 4 for 500hours
Figure 7.4 Optical images for some materials exposed to 4 different gas mixture at
400°C for 500 hours.
224 Low Temperature Corrosion Tests in Gas Mixtures
after comparison with un-reacted coated A1006 steel (Figure 4.8e). Examination of the coating thickness also showed no significant change after reaction at such temperature.
In Gas 2, the scales grown on A1006 steel, 2.25Cr1Mo and SG cast iron were very similar to those produced by Gas 1, but slightly thinner. In the case of pure copper
(Figure 7.4r), a two-layered scale consisted of a relatively compact topmost layer over a porous inner layer. Scale-alloy separation was evident. For Cu-S alloy, the inner scale layer was more porous than that formed in Gas 1 after 500 hours exposure (Figure 7.4v) and preferential attack on the sulfide phase was again evident. No observable scale could be found on coated A1006 (Figure 7.4n). Again the coating thickness had not changed after exposure.
In Gas 3, A1006 steel, 2.25Cr1Mo and SG cast iron developed scales similar to, but thinner than those formed in Gases 1 and 2. The scale formed on pure copper was much more compact than those formed in Gases 1 and 2. Individual oxide grains were clearly visible. In the case of Cu-S alloy (Figure 7.4w), non-protective scale is similar to that observed previously in Gases 1 and 2 was formed. For zincalume alloy, the zigzag appearance of the metal surface resulted from preferential attack on the zinc-rich phase when it was exposed to HCl-containing gas. For pure aluminium, no corrosion product could be seen. Again, no significant amount of corrosion products could be found and no significant thickness change could be observed on coated A1006 tested in Gas 3.
In Gas 4, scale structures similar to those previously observed in other gas mixtures formed on A1006 steel, 2.25Cr1Mo and SG cast iron (Figure 7.4d, 7.4h and 7.4l). No sign of iron sulfide formation was visible, although the XRD results showed trace amount of iron sulfide to be present on all these materials. Copper and Cu-S alloy were severely attacked by Gas 4. In the case of pure copper, a two-layered scale consisted of
225 Low Temperature Corrosion Tests in Gas Mixtures
a relatively compact outer layer and a porous inner layer (Figure 7.4t). For Cu-S alloy, the sample was nearly all consumed after 500 hours exposure (Figure 7.4x). Generally, the scale formed on Cu-S alloy was similar to that formed on pure copper. However, its outermost layer was much more porous and fragmented. EDAX analysis showed the dark grey phase inside the scale to be rich in sulfur and copper on both pure copper and
Cu-S alloy (marked in Figure 7.4t and 7.4x). This is consistent with the XRD results which showed trace amount of copper sulfide on both pure copper and Cu-S alloy.
However, it is difficult to distinguish the sulfur-containing corrosion product from those original copper sulfide inclusions which existed inside the Cu-S alloy. The EDAX analysis also showed chlorine enrichment both at the interface of the any two separated layers in the outmost scale region and the front of the innermost layer (marked in Figure
7.4t and 7.4x). No corrosion product could be seen from the cross-sectional view of coated A1006 steel, zincalume or aluminium.
Scale surfaces were examined for some materials exposed to Gas 3 for 500 hours, using FESEM with EDAX. The SEM images are shown in Figure 7.5. In the case of aluminium and zincalume, surface scratches from sample preparation could still be seen. A thin, fragmented oxide scale was visible on both zincalume and pure aluminium. EDAX showed the surfaces of these scales to be rich in oxygen. In the case of 3Cr12, 253MA and SS304, some areas showed signs of being attacked locally.
A netlike corrosion product was seen on the surface of the 3Cr12 scale. EDAX showed it consisted of oxygen, iron and a small trace of chlorine. Similar morphologies but fewer amounts of netlike corrosion products were observed on 253MA and SS304.
Only aluminium and oxygen were detected by EDAX on coated A1006 exposed to the chlorination environment for 500 hours. Corrosion products formed on A1006 can be described as needle-like and nodular in structure.
226 Low Temperature Corrosion Tests in Gas Mixtures
Zincalume
Figure 7.5 SEM images for materials exposed to gas 3 for 500 hours at 400°C.
227 Low Temperature Corrosion Tests in Gas Mixtures
7.4 DISCUSSION
At such low temperature, negligible weight gains were expected for stainless steels including SS304, 253MA and 3Cr12, coated A1006 steel, aluminium and zincalume.
This made measurement of reaction kinetics difficult. The occurrence of partial scale spallation also strongly affected the observed weight changes of those materials and created further difficulty in analysing their corrosion kinetics.
Results obtained for reaction with air at 400°C were as expected, XRD results showed hematite and magnetite formed on A1006 steel, 2.25Cr1Mo and SG cast iron.
Metallographic study of these materials showed that the thickness of magnetite was thicker than that of hematite. This is consistent with other investigators [10-12] who observed one-stage parabolic kinetics for the oxidation of iron at such temperature. It was explained that the growth of hematite is the controlling step at 400°C. For 3Cr12,
304 stainless steel and 253MA, at such low temperature, XRD analysis could not detect any corrosion products due to their negligible amount. In the absence of more precise measurements, the kinetics cannot be evaluated. Zincalume, pure aluminium and aluminised A1006 steel were found to form very small amounts of corrosion product.
This is expected due to the relatively stable slow growing alumina formed on these materials. For copper and Cu-S alloy, the scale constitution is in good agreement with other studies [26-29].
For all materials tested, the effect of the addition of water vapour was small, as shown by the negligible difference in weight gains and the formation of the same corrosion products.
Reaction with the gas containing hydrogen chloride without sulfur dioxide, led to the formation principally of oxide as shown by XRD results and metallography. Only
228 Low Temperature Corrosion Tests in Gas Mixtures
in the case of pure copper and Cu-S ally was a chloride (CuCl) detected by XRD, although traces of chlorine were found by EDAX in the scale surface developed on the stainless steels. The reaction product identifications are broadly consistent with the observed corrosion rates. Thus the presence of HCl in the oxidising gas had negligible affect on reaction kinetics in the case of low carbon steel, low alloy and stainless steel,
SG cast iron and aluminised A1006 steel. It is therefore concluded that the prediction of the thermochemical diagram in Figure 7.1a are largely borne out: oxides are stable with respect to chlorides for these materials.
The small amounts of chlorine found in the thin scale present on stainless steels shows that this conclusion is an approximation. Identification of the chloride was not possible. However, reference to Figure 7.1a shows that NiCl2 is the most stable of the available chlorides. Whichever chloride formed, its presence indicates that a chlorine- bearing species was able to penetrate the oxide scale to a depth where the oxygen potential was lowered to a level permitting metal chloride formation. The diffusion path representing the growth of such a phase assemblage is shown schematically in
Figure 7.6.
It is seen from Figure 7.1a that the metals for which the chloride-oxide equilibrium is closest to the bulk gas composition are, in addition to nickel, zinc and copper.
Reaction of zincalume with Gas 3 led to net weight loss at all reaction time, consistent with the formation of and volatilisation of ZnCl2 (L). Its vapour pressure together with those of other species is seen in Table 7.3. The zincalume cast alloy is two-phase solid solution–fcc (Al) denoted as (α Al) (light grey phase) or (α′Al) (dark grey phase)
(Figure 7.4y and 7.4z) on the Al-rich and Zn-rich sides, respectively. The zinc-rich
229 Low Temperature Corrosion Tests in Gas Mixtures
phase would be expected to undergo preferential attack, as was in fact seen in Figure
7.4y.
Figure 7.6 Schematic diffusion path for formation of a scale consisting of chloride
and oxide underneath single-phase oxide.
*: Dashed and doted lines are the calculated diffusion path and actual
diffusion path, respectively.
As seen in Figure 7.1b, Cu2O and CuO were expected to form under the reaction gases. However, both pure copper and Cu-S alloy sustained accelerated corrosion and weight gains when exposed to Gas 3. Obviously, the addition of HCl had strong effects on the corrosion mechanism. To trigger the chlorination reaction, the partial pressure of chlorine ( P ) must exceed the required level. Therefore, the inward diffusion of a Cl 2 chlorine containing species is essential. Although difficulties exist in calculating the
P values inside the oxide layer, it is still possible to calculate values for some limited Cl 2 cases (at the oxide layer boundaries). The calculation is based on the following assumptions:
1. No H2O from the bulk gas can diffuse into the scale, and
230 Low Temperature Corrosion Tests in Gas Mixtures
2. HCl can diffuse through the scale freely.
Setting P as the same value in the bulk gas and P as the calculated values at the HCl O2 sublayer boundaries, respectively, P values at the CuO/Cu O and Cu O/Cu interface Cl2 2 2 were calculated using Deacon Reaction [217] (Eq. 7.1) and are shown in Table 7.4. The calculated diffusion path together with the actual diffusion path is shown in Figure 7.6.
Their difference indicates that the simplifying assumptions employed in the calculation are not quantitatively correct.
Table 7.3 Vapour pressures of metal chlorides at 400°C
Species Vapour Pressure (atm)
FeCl3(g) 1.8E-02
CrCl3(g) 1.3E-07
NiCl2(g) 1.4E-08
CuCl(g) 1.0E-10
ZnCl2(g) 7.9E-04
AlCl3(g) 5.1E-01
The formation of solid CuCl, detected by XRD in both cases can be explained using equation 7.3.
2HCl + 0.5O2 = H2O + Cl2 7.1
231 Low Temperature Corrosion Tests in Gas Mixtures
Cu + 0.5Cl2 = CuCl 7.2
Cu2O + Cl2 = 2CuCl + 0.5O2 7.3
MCl2 (g) + 0.5O2 = MO + Cl2 7.4
Table 7.4 Calculated equilibrium Cl2 potential at layer boundaries in copper
oxide scale at 400°C.
Interface P (atm) Cl 2
Cu/Cu2O 2.9E-09
CuO/Cu2O 4.5E-07
-1 At 400°C, CuCl is much more stable than Cu2O (Eq.7.3, ∆Gf = -18676 Cal mol ).
At the low aO values prevailing within the scale, aCl can be high enough to favour reaction (7.3). Then the solid CuCl will develop underneath the previously formed
Cu2O layer. Subsequently, the partial pressure of oxygen will be raised due to the consumption of chlorine and a new Cu2O layer is formed again. As a result, more cracks and scale layer separation would be expected for the corrosion products formed in these two gas mixtures containing HCl due to the volume changes accompanying the oxide formation and conversion to chloride. A schematic drawing of this structure is shown in Figure 7.7. Obviously, the chlorination mechanism observed in the above processes is a modification of the traditional “activated oxidation” mechanism [101], which suggested that gaseous chloride is formed at where the partial pressure of oxygen is low volatile metal chloride diffuses outward and is oxidised and chlorine (Eq. 7.4)
232 Low Temperature Corrosion Tests in Gas Mixtures
could also be further released to penetrate through the scale and react with fresh metal underneath scale formed previously. The “activation oxidation mechanism” was schematically explained in Figure 2.14. As the result of the “activated oxidation”, the corrosion products are very porous and no chloride would be detected. The situation with copper is different, because of the lower copper chloride volatility and the consequent retention of solid chloride in the scale. Unfortunately, it is not possible to compare the results with others, since information related to the chlorination of copper is still lacking. Copper oxides layers CuCl
Gap
Cavities
Cu
Figure 7.7 Schematic drawing of the scale structure of pure copper and Cu-S alloy formed in gas 3 and gas 4 at 400°C.
In Gas 4, the addition of SO2 led to faster corrosion rates for low carbon steel, low alloy steel, SG cast iron, and stainless steels compared to their behaviour in Gas 3. In the case of 304 stainless steel and 253MA, XRD analysis identified only their metal substrate due to the negligible amount of corrosion products formed. For 3Cr12 steel, the corrosion products formed were similar to those of 2.25Cr1Mo steel. XRD analysis on pure aluminium and zincalume showed results similar to those obtained in the other three gas mixtures. According to the prediction shown in Figure 7.1b and 7.1c, all materials tested should be stable in the form of sulphate. However, XRD result showed only metal oxide and trace amounts of metal sulfide were formed. As described in the
233 Low Temperature Corrosion Tests in Gas Mixtures
high temperature sulfidation/oxidation experiments, this indicates that the equilibrium between bulk gas and metal at the metal surface was not reached. A possible reaction involved is shown below:
5Fe + 2SO2 = Fe3O4 + 2FeS 7.5
This has been suggested before [3, 208-212] in explaining iron sulfide formation.
Metallography study showed the corrosion products formed on these materials were similar to those formed in other gas mixtures. Because of the high oxygen partial pressure, the corrosion products were mainly metal oxide. However, metal sulfides may be formed where the oxygen potential is low. In the case of copper and Cu-S alloy, the simultaneous growth of Cu2O and Cu2S may occur:
6Cu + SO2 = 2Cu2O + Cu2S 7.6
This assumption was proved by EDAX result which showed copper sulfide formed at the front of the inner scale layer. The formation of copper sulfide further may have supplied a more ready path for chlorine to penetrate through the scale and therefore the total corrosion rates for both materials were increased. The chlorination of these two materials was explained previously. In the case of zincalume, aluminium and aluminised A1006, no obvious difference could be found from the case of the reaction with Gas 3. The effect of suffer was negligible presumably because of the greater stability of the aluminium oxide scales developed by these alloys.
7.5 Conclusion
i Materials were tested in different gas mixtures at 400°C. The weight changes
for materials that have better corrosion resistance (such as coated A1006,
253MA, SS304, pure Al, zincalume and 3Cr12) were negligible. Difficulties
234 Low Temperature Corrosion Tests in Gas Mixtures
were encountered in measuring corrosion kinetics due to the irregular weight
changes and partial scale spallation. ii Copper and Cu-S alloy were found not to be suitable in gas containing
hydrogen chloride and sulfur dioxide. iii In a sulfur dioxide containing gas mixture, the equilibrium at the interface of
bulk gas and metal surface was not reached. Sulfur dioxide was believed to be
the main reactant that caused the formation of metal sulfide. iv The materials with superior corrosion resistance were: 253MA, SS304 and
coated A1006 steel. v The aluminide coating on A1006 steel showed no signs of deterioration after
these exposure tests.
235 Summary and Conclusions
CHAPTER EIGHT
Summary and Conclusions
A diffusion coating procedure was successfully developed for a mild steel, A1006, by applying the pack aluminization technique. It was found that the coatings grown on
A1006 were critically depended on the activator type. The volatile type NH4Cl activator was found unable to sustain the partial pressure of the aluminium transporting species at the thermodynamically calculated values. This was due to the loss of aluminium transporting species from the semi-sealed crucible. Under those conditions, only FeAl and FeAl 3 phases were formed on A1006 steel, with a relatively thin coating thickness. However, using condensed NaF activator, a 160 mm thick coating with higher rich iron aluminides (Fe2Al5, FeAl 2, as well as FeAl and Fe3Al) was formed after
8 hours. Other factors that may affect the coating growth (such as temperature, masteralloy content and packing density etc.) were not investigated in the present study since they were not the major interests.
It was found that after applying yttria slurry onto the steel surface before the heat treatment, an alumina inclusion free coating was obtained. The formation of the coating was mainly controlled by the outwards diffusion of iron. The kinetics of coating growth in all cases were found to be parabolic.
Coated A1006, together with other materials were evaluated in simulated aluminium smelting cell environments. Different experiment conditions were designed to approximate to different locations above an aluminium smelting cell.
236 Summary and Conclusions
At 800°C, aluminized and bare A1006 steel were exposed to air together with SG cast iron and chromium-containing steels to evaluate their isothermal and cyclic oxidation resistance. TGA was used to obtain their isothermal oxidation kinetics. It was found that weight gain kinetics were parabolic except for SG cast iron. The unusual kinetics observed on SG cast iron reflected a total weight loss in the initial stage followed by a slow weight increase. It is believed that the unusual kinetics were caused by the combined effects of weight loss due to the consumption of graphite exposed to air and the accumulation of iron oxides. Better oxidation resistance was found for
3Cr12, SS304 and 253MA, and attributed to the formation of a protective Cr2O3 layer.
In long term isothermal oxidation experiments, partial scale spallation was found on
253MA and caused the difficulty in analysing the kinetics.
Aluminized A1006 steel was found to have superior oxidation resistance in both
TGA and long time isothermal exposure experiments. Its oxidation rate was found to be very close to that of SS304 stainless steel. In cyclic oxidation experiments, SS304 suffered scale spallation and is therefore considered to be unsuitable for cyclic oxidation conditions. Slow oxidation rates were observed on 253MA and aluminized A1006, and their better oxidation resistance were caused by the formation of protective chromia and a-Al2O3 scale, respectively. These two materials were found to be the top ranked for superior oxidation resistance among all the materials.
The results of tests in sulfidation/oxidation experiments at 800°C showed that equilibrium at the interface between bulk gas mixture and metal was not reached in the current experiment condition. As a result, a duplex corrosion product was observed on bare A1006 steel, 2.25Cr1Mo steel and SG cast iron. Their rapid corrosion was caused mainly by the formation of fast growing iron sulfide phase. The SG cast iron showed
237 Summary and Conclusions
somewhat better sulfidation resistance, which can be attributed to the blockage effect of graphite incorporated into the scale. Alloys 3Cr12 and SS304 were found to perform poorly in the sulfidation experiments, due to the formation of the fast growing iron sulfide. Again, aluminized A1006 steel and 253MA were shown to have the best sulfidation resistance among all the materials tested. This was a result of the formation of protective chromia and a-Al2O3 scale, respectively. Scale formation on these materials was controlled by both outwards diffusion of metal and inwards diffusion of oxidants. Reaction kinetics were consequently parabolic.
Difficulties were encountered in analysing the kinetics of low temperature corrosion reactions, since only small amounts of corrosion products were formed on most materials. Partial scale spallation also occurred in some cases, and resulted in irregular weight changes. Consequently, it was difficult to determine the effect of different gas species on these tested materials. Copper and Cu-S alloy showed unacceptable weight increases when hydrogen chloride gas was involved. However, the presence of the solid CuCl phase inside the copper oxide scale indicates the chlorination of these two materials are different from the traditional “activation oxidation mechanism” [101]. It should be pointed that fluorine containing species are the major reactant inside the aluminium smelting cell. If fluorine were present at the same levels as chlorine used in the present study, the corrosion products would include solid state fluorides. Obviously, it is difficult to predict the scale behaviour of those fluorides at this stage since detailed information of those fluorides is still lacking. Further investigation of the effect of HF on steel structure through other methods would be highly desired. In sulfur dioxide containing gas mixture, equilibrium at the interface between bulk gas mixture and metal was not reached. Sulfur dioxide was believed to be
238 Summary and Conclusions
the main reactant. The best three materials for corrosion resistance were: SS304,
253MA and coated A1006 steel.
Overall, 253MA and coated A1006 steel were the top ranked materials under all conditions. SS304 was found to be unsuitable in high temperature cyclic oxidation and high temperature sulfidation. The 3Cr12 steel also suffered fast corrosion in sulfidation experiments. However, SS304 and 3Cr12 steel showed better performance in low temperature corrosion tests. Alternately, zincalume and pure aluminium can also be considered as potential substitutional materials at low temperature. SG cast iron was found to have better corrosion resistance than to bare A1006 steel and 2.25Cr1Mo steel in all the conditions tested.
In summary, pack aluminization was shown to produce an aluminium-rich iron aluminide coating on low carbon steel. This coating had superior corrosion resistance in several highly corrosive environments, equivalent to that of a high cost stainless steel.
The economic potential of this route to corrosion resistance seems to be worth investigating.
239 References
CHAPTER NINE
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257 APPENDIX 1: LIST OF DATA USED IN COMPUTER CALCULATIONS (Temperature: 1173K Pressure: 1 atm)
Table A.1: Table of the Gibbs Free Energy used in NH4Cl pack.
Species Gibbs Free Energy (Cal/mole) Ar (g) 0 Al (g) 42404 AlC l (g) -34651 AlCl2 (g) -66379 AlC l3 (g) -124759 Al2Cl6 (g) -239579 Al2O (g) -67969 Al2O2 (g) -111409 AlO2 (g) -46052 Fe (g) 57641 FeCl (g) 7543 FeCl2 (g) -46230 FeCl3 (g) -54822 H2O (g) -43734 O2 (g) 0 Cl2 (g) 0 H2 (g) 0 NH3 (g) 19606 HCl (g) -24333 AlN -45726 Fe 0 Al2O3 -311306
258
Table A.2: Table of the Gibbs Free Energy used in NaF pack.
Species Gibbs Free Energy (Cal/mole) Ar (g) 0 Al (g) 42404 AlF (g) -85059 AlF2 (g) -172198 AlF 3 (g) -271585 Al2F6 -550079 Al2O (g) -67969 Al2O2 (g) -111409 AlO 2 (g) -46052 FeF (g) -17676 FeF2 (g) -104530 FeF3 (g) -187073 Na (g) 0 NaF (g) -84269 Na2F2 (g) -190203 Fe (g) 57641 F2 (g) 0 H2O (g) -43734 O2 (g) 0 H2 (g) 0 NaF (s) -108968 Fe (s) 0 Al2O3 (s) -311306
259 APPENDIX 2: XRD RESULTS FOR LONG TERM ISOTHERMAL OXIDATION
EXPERIMENTS.
a 3 Figure A.
260
Figure A.3b
261
Figure A.3c
262
Figure A.3d
263
Figure A.3e
264
Figure A.3f
265
Figure A.3g
266
Figure A.3h
267
Figure A.3i
Figure A.1 XRD results for SS304, 253MA and coated A1006 steel oxidized in air at
800°C after different exposure times.
(a) – (c): SS304; (d) – (f): 253MA; (g) – (i): Coated A1006 steel.
268