HIGH TEMPERATURE SULFIDATION

BEHAVIOUR OF AN Fe-Mn-Al ALLOY

A Thesis submitted in fulfillment

of the requirements for admission

to the Degree of Master of Science

by

Noreen Sa Na Quan

December, 1984

School of Chemical Engineering and Industrial Chemistry

The University of New South Wales STATEMENT

This is to certify that the work presented in this thesis was carried out in the School of Chemical Engineering and

Industrial Chemistry, The University of New South Wales, and has not been submitted previously to any other university or technical institution for a degree or award.

Noreen Sa Na Quan (i)

ABSTRACT

The aim of this research was to investigate the sulfidation behaviour of a ferrous alloy containing 4.5 w/o (weight

percent) manganese, 8.8 w/o aluminium and 0.36 w/o carbon. The alloy was sulfidised at 973 K, 1073 K and 1173 K in H2/H2S atmospheres corresponding to sulfur partial pressures

between 10 -8 atm and 10 -4 atm. The weight gain of each sample over a period of time was measured gravimetrically. X-ray diffraction analyses, microscopic techniques and electron probe microanalysis were used to analyse the scales formed.

Two main types of reaction kinetics were observed. On the one hand, slow parabolic weight uptake kinetics were observed when the combination of temperature and sulfur partial pressure was _ i not too severe at T = 973 K when P <10 atm and at T = b2 1073 K when P = 10 -7 atm. Under these conditions, a thin b2 uniform reaction product layer of oc - MnS and in one instance Al2S3 as well was formed. When only

On the other hand, at higher temperatures and sulfur partial -5 -4 pressures i.e. at T = 973 K when P = 10 atm, 10 atm and _2 at T = 1073 K when 10 < P_ < 10 4 atm and at T = 1173 K b2 when 10 < P <10 atm, breakdown of the protective S2 oc - MnS scale was observed with the formation of large sulfide nodules. This was associated with initially accelerating reaction followed by fast linear kinetics. The (ii)

large nodules were found to contain Fe(Mn)S and a voluminous, porous two phase reaction product layer at their base which was found to be Fe(Mn)Al2S4 plus Mn(Fe)S. This pattern of behaviour was ascribed to mechanical failure of the initially formed protective scales allowing gas access to the under­ lying metal which leads to nucleation and growth of nodules.

Behaviour of a somewhat intermediate type was observed at T = 973 K and P = 10 ^ atm. z (iii)

LIST OF TABLES

Page

Table 1 Mean Rate Constants obtained during 35 the sulfidation of Fe-Al alloys and

Pure Fe. Table 2 Rate Constants obtained during sulfidation 51

of Fe-Mn alloys, Pure Iron and Pure

Manganese. Table 3 Values of the Parabolic Rate Constant, 53 k , for growth of oc - MnS on Mn. Table 4 Sulfidation Rate Constants of 63 Fe-25Mn-5Al alloy. Table 5 Parabolic Rate Constants of 63 Fe-25Mn-20Al.

Table 6 Comparison of Sulfidation Rate 73 Constants at 973 K. Table 7 Comparison of Sulfidation Rate 74

Constants at 1073 K.

Table 8 Standard Free Energies of sulfide 77

formation. Table 9 Equilibrium data for sulfide formation. 79

Table 10 Experimental conditions. 80

Table 11 Sulfidation Rate Constants of 94

Fe-4.5Mn-8.8A1-0.36C. Table 12 Parabolic Rate Constants of Iron and 138 Fe-4.5Mn-8.8A1-0.36C alloy. (iv)

Page

Table 13 Parabolic Rate Constants for 142 sulfidation at T = 973 K.

Table 14 Parabolic Rate Constants at 142 T = 1073 K and P = 10_7 atm. z Table 15: Linear Rate Constants obtained by 149 sulfidation of Fe-4.5Mn-8.8Al and

Fe-25Mn-5Al. (V)

LIST OF FIGURES

Page

Figure 1 Activity gradients of both metal 6 and oxygen established across the

scale. Figure 2 Diffusion of metal through scale. 7 Figure 3 Ellingham Diagram 14

Figure 4 FeS phase diagram in the Fe^_^S 19

region. Figure 5 Autoradiogram of cross-section of 30 sulfide scale on iron at 973 K. Figure 6 Reaction Mechanism (schematic). 38 Figure 7 Manganese concentration vs lattice 44

parameter of (X -manganous sulfide in 1 atm of sulfur vapour at 1073 K, 1173 K, 1273 K and 1373 K. Figure 8 Change in the lattice parameters 44 with sulfur pressures of 1 atm, 10 atm, 10 ^ atm, 3.5 x 10 ^ atm, all at 1273 K.

Figure 9 Lattice parameters (a0,c0) of ferrous 45

sulfide at various pressures at 1273 K.

Figure 10 Equilibrium phase diagram of an 45 FeS-MnS system.

Figure 11 Changes in (Ns/Npe + Nm^) in sulfides 46 with manganese content at various

sulfur pressures at 1273 K. (vi)

Page

Figure 12 Arrhenius plots of the self diffusion 61 coefficient of Mn in the <*: - MnS scales. Figure 13 Schematic assembly of sulfidation 82

rig used. Figure 14 Linear plots, weight gain vs time, 89 for the sulfidation of the alloy at

973 K in 10_8^ Pq ^ 10~4 atm. b2 -4 Figure 15 Parabolic plots, mg2 - cm vs time 90 (min) for the sulfidation of the alloy

— 8 at 973 K in P = 10 atm and 10 -5 atm. 2 -4 Figure 16 Parabolic plots, mg2 cm vs time (min) 91 for the sulfidation of the alloy at 973 K and P = 10 ^ atm. b2 -4 Figure 17 Parabolic plots, mg2 cm vs time (min) 92 for the sulfidation of the alloy at 973 K and P = 10 6 atm. b2 Figure 18 Linear plots, weight gain vs time, for 95 the sulfidation of the alloy at 1073 K

in 10 ^ P 10 4 atm.

Figure 19 Linear plot , weight gain vs time, for 96 the sulfidation of the alloy at 1073 K

in P = 10 ^ atm. b2 -4 Figure 20 Parabolic plots, mg2 cm vs time (min) 97 for the sulfidation of the alloy at

1073 K in P = 10 atm. b2 (vii)

Page

Figure 21 Linear plots, weight gain vs time, 98

for the sulfidation of the alloy at

1173 K in 10"7^T P ^ 10-4 atm. b2 -1 Figure 22 Plots of - lnk^ vs 1/T (k ) for the 101 -4 alloy sulfidised at P = 10 atm. b2 -1 Figure 23 Plots of -lnk^ vs 1/T (k ) for the 102 -5 alloy sulfidised at P = 10 atm. b2 Figure 24 Plots of -log k, vs log (P ) for the 104 b2 alloy sulfidised at 1173 K and 1073 K.

Figure 25 Micrograph of scale formed on alloy 106

sulfidised at T = 973 K and P = 10 ^ b2 atm for 7.4 hours.

Figure 26 Micrograph of scale formed on edge of 107

specimen sulfidised at T = 973 K and -7 P = 10 atm for 4.3 hours. b2 Figure 27 Micrograph inner layer of scale formed 107

on edge of alloy sulfidised at T =

973 K and P = 10 7 atm for 4.3 hours. b2 Figure 28 Micrograph of outermost scale formed 107

on edge of specimen sulfidised at

T = 973 K and P = 10 7 atm for 4.3 b2 hours.

Figure 29 Electron probe microanalysis of 108

specimen sulfidised at T = 973 K and -7 10 atm. (Viii)

Page

Figure 30 Electron probe microanalysis of 109

specimen sulfidised at T = 973 K -7 and P = 10 atm. b2 Figure 31 Micrograph of scale formed on alloy 112

sulfidised at T = 973 K and P = b2 10~6 atm for 4.6 hours.

Figure 32 Micrograph of scale formed on alloy 112

sulfidised at T = 973 K and P = -6 2 10 atm for 4.6 hours.

Figure 33 Tiny nodule formed on alloy sulfidised 112

at T = 973 K and P^ = 10 ^ atm for

4.6 hours.

Figure 34 Electron microprobe analysis of scale 113

formed on nodule of specimen sulfidised

at 973 K and P = 10 ^ atm. b2 Figure 35 Electron microprobe analysis of scale 114

formed on flat surface of specimen

sulfidised at 973 K and P = 10 ^ atm. b2 Figure 36 Micrograph of scale formed on edge of 116

specimen sulfidised at T = 973 K and -5 P =10 atm for 4.8 hours. b2 Figure 37 Micrograph of scale formed on edge of 116

specimen sulfidised at T = 973 K and -5 P =10 atm for 4.8 hours. b 2 Figure 38 Micrograph of scale formed on flat 116

surface of specimen sulfidised at T =

973 K and P = 10 ^ atm for 4.8 hours. b2 (ix)

Page

Figure 39 Micrograph of scale formed on flat 116

surface of specimen sulfidised at

T = 973 K and P = 10 atm for 4.8 b2 hours.

Figure 40 Electron microprobe analysis of scale 117

formed on flat surface of specimen

sulfidised at 973 K and P = 10 atm. b2 Figure 41 Micrograph of scale formed on edge of 119

specimen sulfidised at T = 973 K and -4 P = 10 atm for 4.9 hours, z Figure 42 Micrograph of scale formed on specimen 119

sulfidised at T = 973 K and P = 10 ^ b2 atm for 4.9 hours.

Figure 43 Micrograph of scale formed on alloy 119

sulfidised at T = 973 K and P = 10 z atm for 4.9 hours.

Figure 44 Micrograph of scale formed on alloy 119

sulfidised at T = 973 K and P = 10 z atm for 4.9 hours.

Figure 45 Micrograph of scale formed on specimen 120

sulfidised at T = 1073 K and P = 10 z for 5.4 hours.

Figure 46 Micrograph of scale formed on edge of 122

specimen sulfidised at T = 1073 K and -5 P = 10 atm for 5.9 hours. b2 Figure 47 Micrograph of scale formed on edge of 122

specimen sulfidised at T = 1073 K and

Pc = 10 atm for 5.9 hours. b2 (X)

Page

Figure 48 Micrograph of internal layer of nodule 122

formed on edge of specimen sulfidised

at T = 1073 K and P = 10 ^ atm for b2 5.9 hours.

Figure 49 Micrograph of tiny nodule formed on 123

flat surface of specimen sulfidised at

T = 1073 K and P = 10 ^ atm for 7.8 b2 hours.

Figure 50 Micrograph of tiny nodules formed on 123

flat surface of specimen sulfidised at

T = 1073 K and P = 10 ^ atm for 7.8 b2 hours.

Figure 51 Micrograph of nodule formed on flat 124

surface of specimen sulfidised at

T = 1073 K and P = 10 4 atm for 1.9 b2 hours.

Figure 52 Micrograph of nodule adjacent to 124

protective layer of scales on flat

surface of specimen sulfidised at

T = 1073 K and P = 10 4 atm for 1.9 b2 hours.

Figure 53 Micrograph of specimen sulfidised at 124

T = 1073 K and P = 10 4 atm for 1.9 b2 hours.

Figure 54 Electron microprobe analysis of a nodule 126

formed on specimen sulfidised at 1073 K -4 and P = 10 atm. b2 (Xi)

Page

Figure 55 Micrograph of scale formed on edge of 127

specimen sulfidised at T = 1173 K and -7 P = 10 atm for 4.0 hours. b2 Figure 56 Micrograph of needle-like protrusions 127

into alloy sulfidised at T = 1173 K -7 and P = 10 atm for 4.0 hours. b2 Figure 57 Micrograph of internal layer of scale 127

formed on edge of specimen at T = 1173 K -7 and P = 10 atm for 4.0 hours. b2 Figure 58 Micrograph of scale formed on flat 128

surface of specimen sulfidised at

T = 1173 K and P = 10 ^ atm for 4.0 b2 hours.

Figure 59 Micrograph of scale formed on specimen 128

sulfidised at T = 1173 K and P = 10 2 atm for 4.0 hours.

Figure 60 Micrograph of internal layer of scale 129

formed on specimen sulfidised at T =

1173 K and P = 10 6 atm for 3.9 b2 hours.

Figure 61 Micrograph of external layer of scale 129

formed on specimen sulfidised at T =

1173 K and P = 10 ^ atm for 3.9 b2 hours.

Figure 62 Micrograph of scale formed on alloy 130

sulfidised at T = 1173 K and P = O ^ atm for 2.7 hours. (Xii)

Page

Figure 63 Micrograph of scale formed on 130

specimen sulfidised at T = 1173 K - 5 and p = 10 atm for 2.7 hours. b2 Figure 64 Micrograph of scale formed on 130

specimen sulfidised at T = 1173 K -5 and P = 10 atm for 2.7 hours. b2 Figure 65 Micrograph of scale formed on edge 131

of specimen sulfidised at T = 1173 K -4 and P = 10 atm for 1 hour. b2 Figure 66 Micrograph of external layer of 131

scale formed on specimen sulfidised

at T = 1173 K and P = 10~4 atm b2 for 1 hour.

Figure 67 Micrograph of internal layer of 131

scale formed on specimen sulfidised

at T = 1173 K and P = 10 4 atm for U rj

Figure 68 Electron microprobe analysis of edge 133

of specimen sulfidised at 1173 K and

P = 10"4 atm. b2 Figure 69 Activity gradients of Mn and Al. 141

Figure 70 Proposed Mechanism of Nodule Growth 146 (Xiii)

ACKNOWLEDGEMENTS

I wish to thank Associate Professor David John Young, who as supervisor of this project, gave generously of his time to provide encouragement, technical advice and constructive criticism. Thanks are due also to the staff of the School of Chemical Engineering and Industrial Chemistry at the

University of N.S.W. for their invaluable assistance.

Mrs. Judy van Raak is especially thanked for cheerfully undertaking the extensive task of typing this thesis.

I would also like to thank my husband for his encouragement. (xiv)

TABLE OF CONTENTS

Page Abstract (i)

List of Tables (iii) List of Figures (v)

Acknowledgements (xiii)

Table of Contents (xiv)

1.0 Introduction 1

PART A - LITERATURE SURVEY 2.0 General Sulfidation Theories

2.1 Wagner's Oxidation Theory 5 2.2 The Parabolic Rate Law 10

2.3 Sulfidation of Binary Alloys 15 2.4 Sulfidation of Ternary and Higher 18 Alloys 3.0 The Sulfidation of Pure Iron

3.1 Description of Sulfide Scale 20 3.2 Kinetics 21 3.3 Temperature Dependence of Sulfidation 25 3.4 Pressure Dependence of Sulfidation 26

3.5 Duplex Scale Formation 28 4.0 The Sulfidation of Binary Alloys

Containing Aluminium or Manganese

4.1 Sulfidation of Fe-Al Alloys 33 4.1.1 Kinetics 33

4.1.2 Description of Sulfide Scales 39

and Their Formation Mechanism (xv)

Page 4.2 Sulfidation of Fe-Mn Alloys

4.2.1 The FeS-MnS-S System 47 4.2.2 Kinetics 49

4.2.3 Description of Sulfide Scales 54

and Their Formation Mechanism 5.0 The Sulfidation of Iron - Aluminium -

Manganese Alloys 5.1 Kinetics 64 5.2 Description of Sulfide Scale 67 5.3 Comparison of Sulfidation Rates 75

PART B - EXPERIMENTAL WORK

6.0 Experimental Procedure 6.1 Introduction 76 6.2 Experimental Conditions 76 6.3 Sample Preparation 81 6.4 Sulfidation Apparatus 83 6.5 Characterisation and Identification 86 of Scales 6.5.1 X-Ray Diffraction Analysis 86 6.5.2 Metallographical Examination 86

6.5.3 Electron Probe Microanalysis 87 6.6 Analysis of Data 88 7.0 Results

7.1 Kinetics 93 7.1.1 Temperature Dependence of Sulfidation Rate 103 (xvi)

Page

7.1.2 Pressure Dependence of 105

Sulfidation Rate

7.2 Sulfide Scale Morphology and 110

Phase Constitution

7.2.1 T = 973 K, P = 10_8 atm 110 b2 7.2.2 T = 973 K, P = 10~7 atm 110

b**2 _£ 7.2.3 T = 973 K, P =10 atm 111 b2 7.2.4 T = 973 K, P = 10~5 atm 118 b2 7.2.5 T = 973 K, PQ = 10~4 atm 121 b2 7.2.6 T = 1073 K, P = 10_7 atm 121 b2 7.2.7 T = 1073 K, P = 10~6 atm and 125 2 -5 PQ = 10 atm b2 7.2.8 T = 1073 K, P = 10_4 atm 125 b2 7.2.9 T = 1173 K 132

8.0 Discussion

8.1 Introduction 136

8.2 Protective Scaling Kinetics 137

8.3. Fast Nodule Growth Kinetics 147

9.0 Conclusion 152

10.0 References 154

PART C - APPENDIX 155 1

1.00 INTRODUCTION

Iron based alloys are used frequently where high temperature corrosion resistance is required. In particular, they are used by the petroleum industries which have progressed from simple, low temperature distillation -to high temperature, high pressure catalytic cracking and reforming in an attempt to realise greater efficiencies. During the processing of sulfur containing feedstocks, hydrogen sulfides are released.

These cause corrosion and scaling problems in the processing vessel. To prevent this attack at the temperatures of operation (873 K - 1273 K), the coal-petroleum industries often turn to a variety of stainless steels or to more exotic bonded alloys ^ ^. However, these alloys are quite expensive and therefore it is of great economic interest if cheaper alloys could be developed which would be resistant to hydrogen sulphide attack at temperatures of concern.

In the present work, the effects of alloy additions of manganese and aluminium have been investigated. These elements have obvious economic advantages over chromium. At sulfidation temperatures over 873 K, where the traditional stainless steels perform poorly under reducing conditions, they might be expected to be beneficial to alloy sulfidation resistance.

It is well known that Additions of aluminium to iron based alloys increase their resistance to sulfidation due to the formation of a protective film of aluminium sulfide. Aluminium was selected because of the following points: (2) 2

1. Aluminium sulfide is thermodynamically more stable

than FeS. 2. The sulfidation rate of aluminium is lower than that of

pure iron at low temperatures. 3. The large Pilling-Bedworth ratio of Al2S3.

The effects of aluminium in sulfidising atmospheres have been studied by many investigators. Strafford (2) and (3 4) Nishida ' have found that the addition of aluminium was quite effective in reducing the overall rate of sulfidation compared to pure iron.

Elrefaie and Smeltzer found that the rate of sulf­ idation of pure manganese was less than that of pure iron.

It has been observed that manganese addition to iron reduces the sulfidation rate and that manganese offers more effective corrosion resistance to iron than chromium under certain conditions .

Very few studies have been done on the sulfidation of ferrous alloys containing both aluminium and manganese.

Smeltzer ^ found that under certain conditions, some

Fe - Mn - Al alloys have higher sulfidation resistance than some Fe - Mn alloys.

It is obvious from the above that the additions of aluminium and manganese are effective in developing sulphidation resistance when present at sufficient concentrations. The present study aims to investigate the sulfidation behaviour of an alloy containing both elements at low concentrations. 3

The following literature survey attempts to establish why metal failure occurs under sulfurising atmospheres. It also seeks to outline the routes to alloy development for sulfidation resistance. PART A - LITERATURE SURVEY 4

PART A - LITERATURE SURVEY

This literature survey is divided into four major sections.

The first is concerned with the theories which are

applicable to sulfidation - Wagner's Oxidation Rate Theory

and the derivation of the parabolic rate law, together with

the theory behind binary and ternary alloy sulfidation.

In the second section the sulfidation of pure iron is

outlined.

The third section reviews previous work done on the

sulfidation of binary alloys containing aluminium or manganese.

In the final section, a brief review of the sulfidation of

iron, aluminium and manganese alloys is presented. 5

2.0 GENERAL SULFIDATION THEORIES

2.1 Wagner's Oxidation Theory

Most investigations of the sulfide scale growth on Fe alloys

have produced results consistent with Wagner's Oxidation

Theory.

The mechanism of corrosion depends on the nature of the

scale, whether it is compact or porous. A compact scale

acts as a barrier which separates the metal and the reacting

gas. If sufficient gas is available at the oxide surface,

the rate of reaction at high temperatures will be limited

by solid-state diffusion e.g. lattice, grain boundary, or

short-circuit diffusion, through the compact scale. As the

diffusion distance increases as the scale grows in thickness,

the reaction rate will decrease with time. Compact scales

offer the best protective properties. In most attempts to

improve the corrosion resistance of metals and alloys, the

goal is to improve the protective properties of the

reaction product scale.

Wagner's well known and widely applied theory of high

temperature parabolic oxidation is the most important

single contribution to the understanding of high temp­ erature oxidation of metals. This theory is not limited to oxidation of metals but is also applicable to gas-metal reactions in general.

The main assumptions of Wagner's theory are summarised below:

1. The reaction product layer is a compact, perfectly 6

Figure 1 Activity Gradients of both Metal and Oxygen Established across the Scale.

Metal M Oxide Gas 02

2 + + 2e Cations M + 2e + ^02 = MO

kO 2 = 2e = Q2 _ M + 0 MO + 2e Anions

Electrons

II

2

l exp 2 AG 02 7 M/MO exp (V h

ii p oxygen partial pressure at scale-gas interface o 2

I activity of metal at scale-metal interface 7

Figure 2 Diffusion of Metals through Scale

I Cvm = concentration of cation vacancies

C, = concentration of electron holes 8

adherent scale.

2. Migration of ions or electrons across the scale is the rate controlling process. 3. Thermodynamic equilibrium is established at both the

metal-scale and scale-gas interfaces. 4. The scale shows only small deviations from

stoichiometry. 5. Thermodynamic equilibrium is established locally

throughout the scale. 6. The scale is thick compared with the distance over which space charge effects (electrical double layer) occur. 7. Reactant gas solubility in the metal may be neglected.

Since thermodynamic equilibrium is assumed to be established

at the metal-scale and scale-gas interface, it follows that activity gradients of both metal and non metal (oxygen, sulphur, etc.) are established across the scale as shown diagramatically in figure 1.

Consequently, metal ions and oxygen ions will tend to migrate across the scale in opposite directions. Since the ions are charged, this migration will cause an electric

field to be set up across the scale resulting in consequent transport of electrons across the scale from metal to atmosphere. The relative migration rates of cations, anions and electrons are therefore balanced such that no net transfer occurs across the oxide layer as a

result of ionic migration. 9

Being charged particles, ions will respond to both chemical and electrical gradients which together provide the net driving force for ionic migration.

If the reaction product is defective within the cation sublattice e.g. Fe1_yS, the scale grows as a result of out­ ward diffusion of metal as shown in figure 2. 10

2.2 The Parabolic Rate Law

The mobility or the diffusion coefficients for the cations,

anions and electrons are not equal, and because of this

difference, a separation of charges take place in a growing

scale. But the resulting space charge also creates an

electric field which opposes a further separation of charges

and a stationary state is reached for which no net electric

current flows through the scale (7)

Jz, q J, = 0 ...(1) l l where = valence of species i, q = electron charge

and J = particle flux (mole no. of particles /cm2s)

In describing the transport of ions and electrons through

the scale, both the diffusion due to the chemical potential,

JJL, and that due to the electrical potential, V, must be

taken into account. On this basis, the number of particles

passing through 1 cm2 of a plane per second is given by:

T Ci Di ^V± ...(2) Ji RT £ x

where J. = each individual particles flux, Ch = concen­

tration, 7^i = electrochemical potential and D_^ = diffusion

coefficient.

The electrochemical potential is defined by:

yj = /J- + Z qV ...(3)

where V = electrostatic potential and y(t = chemical

potential and hence:

5v 4 B x + Z q B x . • . ( ) 11

In equation (4) , dM: is the chemical potential gradient,

and - dv is the electric field strength with the dx position co-ordinate, x, being chosen normal to the

surface of the reacting metal.

iron sulfidises to form Fe„ S, a metal deficit scale in 1-y which y is a measure of deviation from stoichiometry. In

this sulfide, J = 0 as the sulfur does not diffuse ' s (44 ) inwards through the scale.

As the sulfur species are immobile, the only particles of interest are cation vacancies and positive holes. The

evaluation of the chemical potentials of these species may be obtained from the following solid-gas equilibrium:

I S2 (g) = + VF; = 2 h ...(5)

(77) Here the notation of Kroger and Vink is employed.

In this notation the principal symbol indicates the species

in question, the subscript denotes the site on which the

species resides, and the superscript shows the charge of the species relative to normal site occupancy, thus

X Sg = sulfur anion at a sulfur position in the lattice

with no effective charge

I! VFe = an Fe ionic vacancy with a relative charge of 2 -, h = a positive hole with a 1 + charge

Since the activity of the lattice sulfide ions is essentially constant,

dy^s = a Uv + 2dyth . ..(6) 12

Because the vacancies are double charged

d^v + 2 d7|h dy^v + 2 dy^h

d^v + 2 dyjh dyUs

RTdlna = RT d In P, h 2 . (7)

The zero net current condition for Fei_yS is

Jh 2 J, . (8)

(29) Since Fei_yS is a p-type semiconductor

Ch Cv Dv . (9)

Combining equations (2), (8) and (9) gives:

h 2 Cv D v I 2^2 . (10) 2 ch Dh v ^ 5

After some algebra, the following is obtained: P. d In S2^ — C v D v . (ID

d J Now 0 under steady state conditions, and

integrating equation (11) gives:

x = X x — X \ Jv dx = - \ Cv Dv d 1 n P . (12) x = o x = o

X — A J X = - D \ Cv d In Pg '■ . (13) x = o ^ where X is the thickness of the scale. 13

Now from the supposed point defect equilibrium shown by equation (5)

Cv K P. .•(14) whence JX = - Dv K (PQ 6 (x=X) - PQ 6 (x=o) ) b2 b2

Since Vm is the molar volume of vacancies, then:

dX J Vm dt

...(15) X 1 where k' = - Dv K P 6 and P (x=o) is assumed to be S2 S2 negligible.

Now integration of equation (15) gives

I X dX

X2 k' t 2

X2 = 2k' t .•(16)

This is the parabolic rate equation for scale growth. 14 H n S/ Ho Ratio Ps2 (atm) 10 10 10 10 10J 1200 102 10

i CqS. 1 •2FeS 1-2

-3

2Cr S 1 o6

;8 10 2MnS

2 T i S

-160

io14

-200

-2 40

-22 10 Absolute i°t 10 io 1QV 10 zero

Figure 3: Ellingham Diagram 15

2.3 Sulfidation of Binary Alloys

The mechanism of sulfidation of binary alloys is considerably more complex than the mechanism of sulfidation of pure metals. The principles governing the oxidation of binary alloys apply equally well to binary alloy sulfidation and will be summarised in this section. The morphology and phase composition of sulphide scales on alloys depend on alloy composition and on the ability of the sulphides of alloy components to form solid solutions or spinel structures. The various scale morphologies arising from the oxidation and sulfidation of binary alloys have been recently summarised by Bastow, Wood & Whittle1 ' and Young (9)

There are two categories of external sulfide scale - single or multi phase (9) . When the component metals can form sulfides of similar thermodynamic stabilities and crystal structures, substitutional solid solubility is possible resulting in a single-phased mixed sulfide scale. To form a single phase mixed sulfide on a binary alloy, the second metal must be chemically very similar to iron and its diffusion coefficient in the mixed sulfide is therefore expected to be similar to that of iron.

As can be seen from the Ellingham diagram in figure 3, it is rare to find binary alloys whose components form sulfides of similar stability.

A number of the elements normally alloyed with iron form sulfides which are more stable than Fei_yS. In this case, 16

the more stable sulfides will form at the place where the sulfur activity is lowest: either beneath the alloy surface as an internal precipitate or within the scale and next to the alloy surface. If the more stable sulfide is an external scale constituent then two limiting cases arise: 1. When iron is partially soluble in the other sulfide,

a coherant sublayer of that sulfide is formed on the alloy surface and an outer sublayer of Fei_yS is formed by the diffusion of Fe through the sublayer. 2. When iron is negligibly soluble in the other sulfide, no mechanism for the formation of Fei_yS would exist if such a morphology was formed.

Because the formation of Fei_yS is thermodynamically required at sufficiently high Ps2, an alternative morphology is formed in which an inner two phase sublayer of Fe^yS plus the second sulfide is formed under an external layer of Fei-yS. In the first case, if the more stable sulfide has a low diffusivity for iron, a passivating effect is possible. In the second case, the effective cross-section for Fei_yS growth is reduced as the other sulfide has a blocking effect on iron diffusion and some reduction of rate could be expected.

As previously mentioned above, when an alloy additive which forms a more stable sulfide than FeS is added to iron, internal sulfidation is possible. Many metals form more stable sulfides than FeS (see fig. 3) and as alloy additives, are able to form the only sulfide when the 17

sulfur partial pressure is less than that characteristic of the Fe/Fei_yS equilibrium. Internal sulfidation will occur when these alloy additives are dilute enough.

It has been shown that the internal sulfidation of dilute Fe - Mn alloys at 923 - 1123 K follows parabolic kinetics. Wagner's theory was used to estimate the diffusion coefficient of sulfur in Fe, Ds which was an order of magnitude greater than an independent measurement

(11) The metallographic cross sections revealed increased grain boundary penetration of the sulfur.

Wagner (12) showed that when there is no external scale and when the rate controlling step is the inward diffusion of sulfur, then parabolic kinetics are predicted. It has also been shown by Rapp (13) that internal sulfidation can also occur within an alloy which forms an external scale. 18

2.4 The Sulfidation of Ternary and Higher Alloys

As with binary alloys, the phase composition and morphology

of sulphide scales on ternary and multicomponent alloys are very complex. As a rule, the scale is multiphase and has many layers different in morphology and phase composition.

Even though the structure of the scale is complex, their

sulfidation generally follows a parabolic rate law although deviations are sometimes observed. The prediction of the sulfidation behaviour of any alloy is impossible without appropriate phase diagrams and diffusional information. 19

CD ■ SH 2 CD 2 Sh i—i 2 2 rd 0 4-) 2 Sh 2 o 04 Sh cy> CD 2 04 x fd g CD •H CD 42 ■P 4-> 0 fd 42 04 Cn •H u X! 0 2 m CD 4- ) i—I X •H 0 4-> I-- 1 2 O -H Sh Cl o g o 5- 1 0 MH 4-> •H 5h 23 CD 43 CD 42 •H 4-1 Eh i—I 4-1 in ■H 0 CNl 2 1 -1 01 04 2 W O CD •H Sh tT> 2 o CD 2 CN 5h £ o 2 co 42 P4 CO 44 i \—1 2 ro (D Sh • Eh 2 i—i LO o 2 2 rH X 2 4J 2 II >i 2 2 X g o •H 04 4-> \ ■ <—! lU • g 2 s T—1 0 •—' tn 2 u 2 LO -H •H P4 23 2 44 o •H • 2 r—1 2 — 2 u •H O'i 2 0 '—- X 2 2 o 04 0 Sh 2 g 2 0 o CO 43 -H 1 23 0 &» o 2 2 2 2 Eh 2 ■H Sh

• H PS D U M Eh 20

3.0 THE SULFIDATION OF PURE IRON

3.1 Description of Sulfide Scale

Below the FeS eutectic temperature (1261 K), pure iron sulfidises to form a compact, tightly adherent scale in the relatively short term. As shown in figure 4, at sufficiently high P , a thin layer of FeS2 is formed b2 over a thick layer of Fei-yS „ However, at lower P , S2 only Fei_yS is observed.

As the rate of formation of Fei_yS is much faster than

FeS2, attention has been concentrated mostly on the mono­ sulfide formation. 21

3.2 Kinetics

Pure iron has been observed to sulfidise according to parabolic kinetics at sulfur pressures ranging from 10-11 to 1 atmosphere at temperatures in the range 673 K -

1273 K irrespective of whether the FeS scales formed are single layered and compact or duplex, with an inner striated and microporous layer. The parabolic behaviour and the fact that the platinum markers remain at the scale

- metal interface suggest that the rate determining step of the sulfide scale growth on iron is the outward diffusion of metal. It has been shown that the ferrous sulfide is a metal-deficit, mixed electron-hole semiconductor and that the self-diffusion coefficient (15) of iron in this sulfide is higher than that of sulfur The main ionic defects in Fei-yS are therefore cation vacancies. The rate determining step in the growth of a compact sulfide scale on iron is then the outward diffusion of iron ions through cation vacancies in ferrous sulfide.

The mechanism of sulfide scale growth on pure iron can be represented by Wagner's cation vacancy diffusion model i.e. an outward diffusion of iron cations and electrons with a countercurrent inward diffusion of cation vacancies and electron holes v ' ' .As long as the zone of metal consumption can be accommodated by plastic flow of the outward growing scale (i.e. by inward relaxation) the scale remains a compact single phase layer of iron sulfide. 22

It has also been shown that iron sulfide exhibits a wide range of non stoichiometry, up to 25 atomic per cent, and that this is due to the presence of iron vacancies i.e.

Schottky type defects, The degree of non stoichiometry was found to increase with temperature or sulfur partial pressure (19)

Because of the metallic electronic conductivity character­ istics of Fei_yS and because the self diffusion coefficient of sulfur Dg is much smaller than that of iron Dpe, then

Wagner's theory predicts that the flux of iron through a growing sulfide scale is given by Eq. (11) which can be reformulated as:

ZFe d In a, JFe CF« 2 |Z . • • (17)

If the flux is assumed position independent, equation (17) integrates to give the parabolic rate law with the rational rate constant expressed as:

k. DFe d In as .•(18) a' s where is the average concentration, in equivalents, of metal in the scale and a', a" are the sulfur activities at s s the metal-^scale and scale-gas interfaces.

The rational rate constant (in equivalent cm 1 s !) is related to the parabolic rate constant (in g2 cm-4 s-1) by the following equation:

% V k / A ...(19) eq p 23

where V is the equivalent volume of Fe., S, k is the eq ^ 1-y p parabolic rate constant, and A is the atomic weight of S.

Independently measured diffusion data has verified the

quantitative application of equation (18). The differential

form of equation (18):

d k r ...(20) d In a"

for which T and therefore a are fixed, was used by Mrowec s who found order of magnitude agreement between Dpe

deduced from a plot of k Vs In P , and the tracer r b>2 (21) diffusion coefficient of Fe . The dependence of DFg

ZFe ^ Zs °n as Was ne9-)-ecte(^ i-n arriving at equation (20) .

(22) However, it has been shown that the tracer diffusion T coefficient, , in mono crystal Fe., S can be described Fe 2 1-y thus:

Dq y exp [- (81 + 84 y) kJ / RT]

T Therefore DFg and hence D , depends on composition, and

therefore a^, for y small but is nearly constant for larger values. The contention that y and therefore the effective

Fe valence is independent of ag is very questionable although the approximation f (a^) may be acceptable.

(23) Young and Smeltzer evaluated the integral in equation T (18) on the assumption that DFg = DFe and found that agreement between the calculated and measured parabolic rate constants was within half an order of magnitude for 24

806 973 K and 8 x 10 P ( 1 x 10 5 atmospheres. b2^

Diffusion of Fe in Fe„ S is anisotropic with the ratio of l-y the tracer diffusion coefficients in the c and a directions given by (19):

cn 1.8

-3 When the sulfur pressures are less than 10 atmospheres, the scales exhibited a preferred texture with the "a" (19) axis parallel to the direction of growth . At sulfur _ 3 pressures exceeding 10 atmospheres, the sulfide scale grows preferentially along the "c" axis of the hexagonal structure of this sulfide. Consequently, the scale has a well developed texture characterised by columnar crystals situated perpendicularly to the metallic core and parallel to the "c" axis. With this information and the diffusion (19 24) coefficients measured on Fe. S scales, Fryt et.al ' l-y found that the calculated and the measured values of the parabolic rate constants were very close for 873 ^ T ^

1253 K and 5 x 10 P < 2 x 10 2 atmospheres. The b2^ reasons for the change in the preferred orientation have not yet been explained but it is well known that important factors are temperature, P and exposure time. b2 25

3.3 Temperature Dependence of Sulfidation

At constant sulfur pressures, the activation energy of the parabolic rate constant ranges from 0 to 85 kJ/mole according to the value of P . This is due to the strong b2 effect of P on the temperature dependence of y and thus b2 the integral in equation (18). This has been most elegantly shown by Fryt et.al (19) using the Libowitz

(52) point defect model involving a strong repulsive interaction between iron vacancies. Fryt et.al found that the predicted values for the rate constants agreed with those found experimentally over the ranges 10 ^ P <3

The relationship between ag and y is given by Libowitz model as:

In (as/y) = (AFeS + 9V> / RT + 4 Evy (2 - y) /RT where anc* 9V are the standard Gibbs free energies of formation of stoichiometric FeS and iron vacancies, and

is the free energy of interaction between iron vacancies, which is experimentally determined.

This satisfactory explanation of the variation of activation energy with P is further proof of the applicability of b2 Wagner's theory to the sulfidation of iron. 26

3.4 Pressure Dependence of Sulfidation

The sulfur pressure dependence of the parabolic rate

constant is of the form:

1 ...(21) k p c< (PbQ 2 ) n

where n has been reported variously, as having values in

the range 3 to 7.

There is confusion in the literature as to the dependence

of the iron sulfidation rate on temperature and P . b2 Predictions of sulphur partial pressure effects have been

obtained from equation (18) which made use of the

relationship:

D D C .. . (22) Fe V V

To integrate equation (18), as a function of ag must be known and this was obtained from the supposed point

defect equilibrium shown by equation (5) and the charge

balance 2 C = C, . With the further assumption of ideal v h or Henrian solution behaviour, equation (5) gives:

1 1 C = (K/4) 13 ** P* 6 ...(23) v S ~ where K is the equilibrium constant for equation (5).

Combining equation (18), (22) and (5) and integrating produces:

1 6 k r 27

1 However the quantity n has been reported to vary between 1 1 3 & 7 and these results have been interpreted as agreeing with the above point defect diffusion model or, alter­ natively as indicating a degree of gas phase sulfur polymerisation. However, sulfur exists almost entirely as S2 (g) at the T and P of reaction so the polymerisation b2 interpretation is not feasible. The non typical dependence of the sulfidation rate on the temperature and sulfur partial pressure results from the complicated structure of point defects in ferrous sulfide, which is not ideal or Henrian. The observation of a particular pressure dependence is therefore not a good basis for the deduction of a particular point defect mechanism. If, instead, equation (18) is integrated numerically for a number of P values as done by Fryt et.al (24) there is excellent b2 agreement between experiment and theory over the range

10 10<1 Pq ^ 1 atmosphere and 873 T ^ 1253 K. This good b2 agreement gives further support to the applicability of Wagner's description to the iron sulfidation reaction. 28

3.5 Duplex Scale Formation

When the extent of sulfidation of iron is large, a duplex scale is formed which consists of a compact outer layer over a porous layer.

Because of the large size of the sulfur ion, sulfur is not mobile in the Fe1_yS lattice and therefore this morphology cannot be explained by the simultaneous solid-state

2 + 2- diffusion of both reactants: Fe and S ions. When iron is sulfidised extensively, mechanical failure at the scale- metal interface can result in a duplex scale being formed in which the growth of a porous inner layer is supported by inward diffusion of gaseous sulfur while the fairly compact outer layer continues to grow by diffusion of iron.

This generally accepted explanation has been proposed by many workers v(9 ' 25—32') and the mechanism will be expanded below.

The regular parabolic stage of the growth of a compact scale of FeS on iron in the early stages of reaction is controlled 2 + by the outward diffusion of Fe ions from the metal to the scale-sulfur interface. The activation energy for sulfidation agrees with the reported value for Fe diffusion in FeS (33) . The outward diffusion of Fe 2+ ions is, however, accompanied by an inward diffusion of vacancies, which must be accommodated at the Fe surface. When the extent of sulfidation is large, product sulfide is formed at the scale-gas interface where it cannot "fill the gap" left by the metal consumed at the metal-scale interface. 29

If the scale is to remain in contact with the retreating metal surface, it must deform in a plastic fashion. When a limit to the scale's plastic deformation capacity is reached, the scale separates from the metal and voids are formed. This happens easily at sample edges or on convex surface where plastic deformation is not possible due to geometrical constraints. After formation of the voids at the scale-metal interface, the sulfur activity at the scale under-surface rises by the dissociation of the FeS in the region of these voids to yield sulfur vapor, which is able to react with the underlying metal, allowing nucleation and growth of the inner scale. This mechanism of reaction is called the dissociative mechanism (25,27,30,31,32).

Decomposition of the outer layer takes place preferentially along planes perpendicular to the core grain boundaries, (3 0 since these are planes of the slowest crystal growth ' 31,34,35).

Therefore, dissociation of the initial layer of the scale results in microfissures which are perpendicular to the core. At the specimen edges, this process begins almost immediately after the formation of the initial layer which loses its compactness over its entire cross-section.

Further growth of the scale takes place by outward lattice diffusion of ions and electrons of metal and inward diffusion of sulfur molecules through microfissures which also could possibly have resulted from fracture of the outer scale when its plastic deformation limit was being reached. Eventually, the void or gap between metal and the original 30

Figure 5: Autoradiogram of cross-section of sulfide

scale formed on iron at 973 K. Sample

exposed to non-radioactive sulfur for

25 min and then to 35S for 75 min. After

Bruckman and Romanski scale is filled by porous sulfide and outward metal transport starts again resulting in further growth of the outer coarse grained scale.

The effect of surface curvature on the morphological structure of the scales on metals is shown by experiments which consist of initially exposing the metal to non­ radioactive sulfur for a certain period of time and then admitting a radioactive isotope of sulfur S35 into the reactor. As seen in the autoradiogram in figure 5, in the specimen edges, radioactive sulfur is present both in the outer scale, grown in the second stage of reaction, but also in the inner part formed initially with the normal sulfur. This implies that at the specimen corners sulfur has diffused inward through the initially formed scale. However, on the flat surfaces, the inner layer does not contain radioactive sulfur which implies that sulfur has not ingressed through the initially formed scale.

It has been shown * ' ' by the Wagner pellet method that when a growing sulfide layer has no geometric constraints, and when full contact between metal and the product is ensured by applying an appropriate load, then a porous inner layer did not form although the compact layer of sulfide grown reached a thickness of several centimeters.

When the sulfide layer was immobilised, immediately a porous inner layer was formed.

/3 7\ (38) Rickert 1 ’ and Mrowec v ' have used Wagner’s theory of interface instability (39) to explain why the separation of 32

the first formed scale should lead to a porous underlayer, on the assumption that gas phase diffusion across the gap is rate controlling. Irregularities on newly forming sulfide's surface would be magnified under these conditions, eventually resulting in a porous structure.

Bruckman and Remanski ^^ have shown that even trace impurities exert an important effect on the growth mechanism of sulfide scale on iron. It was shown that whereas on spectrally pure iron, the scale in the region of flat surfaces is liable to considerable plastic deformation and produces a tightly adherent single layer phase, on

Armco-type iron (total impurities 0.8%) plastic flow did not occur, and a duplex scale with a porous inner layer was produced. 33

4.0 THE SULFIDATION OF BINARY ALLOYS

CONTAINING ALUMINIUM OR MANGANESE

4.1 The Sulfidation of Fe-Al alloys

4.1.1 Kinetics

A considerable amount of test work has shown that aluminium additions are effective in achieving sulfidation resistance (41) for a ferrous alloy. In 1938, Nauman studied the effects of aluminium additions to iron in pure hydrogen sulfide at various temperatures between 623K to 873K. The corrosion rate was found to go through a maximum with small additions (1-2 wt.%)of aluminium but at 7% aluminium, the corrosion rate was reduced to less than that of carbon steel.

Murakama and Nagasaki (42) reported that even small additions of Al to Fe improved resistance to sulfidation in sulfur vapor, and observed that at 717K and under 1 atmosphere sulfur vapour, 5 wt % Al was enough to suppress the reaction almost completely.

Setterlund and Prescott (43) reported that Fe - 6 wt.% Al is roughly equivalent in resistance to sulfidation to an 18 Cr - 8 Ni stainless steel. They studied the corrosion of Fe-Cr and Fe-Al alloys in H2/H2S atmospheres (total pressure approximately 15 and 35 atmospheres) at a temperature range of 533K to 811K. They found that for atomic percentages greater than 12 percent, aluminium additions were more effective than chromium in passivating 34

an iron alloy surface to sulfur attack.

Strafford and Manifold ' ' ; studied the corrosion behaviour of binary iron-based alloys containing 5, 10 and 20 wt % of aluminium, in S2(g) at pressures of 1 x 10-4 to

1.3 x 10-1 atm. and in H2/H2S atmospheres for the temperature

range, 773 to 1273K. At low temperatures (773K - 973K), the addition of 5% Al reduced the corrosion rate very effectively by a factor of 10 relative to pure iron.

However, at the higher temperatures (1073K - 1273K), this reduction is not maintained. Even increasing the Al in the alloys to 10 and 20 wt % does not significantly lower the corrosion rate.

At 1 atm S2(g), Fe - 5 Al alloy sulfidation kinetics were observed to be linear after an initial short period of exposure during which the kinetics were parabolic. At sulfur pressures lower than 1 atm, either in S2(g) or in

H2/H2S atmospheres, linear kinetics were followed.

At 1173K in H2/H2S atmospheres, Fe -10 Al reacted initially at a gradually accelerating rate followed by a period of slow linear kinetics. At 1273K, the reaction was so rapid

_ 2 (wt gain of 35 mg cm after only 2 minutes of exposure) that it was not possible to obtain accurate weight gain Vs time data.

At 1173K, there was negligible reaction with the Fe - 20%

Al alloy. However, at 1273K the alloy sulfidised according to irregular parabolic kinetics. A parabolic plot of weight uptake data revealed three straight portions. The 35

Table 1: Mean Rate Constants obtained during the

(45,46) sulfidation of Fe-Al alloys and Pure Fe

1. Fe - 5% Al

Temperature kp kl

(K) (g cm-4 sec-1) (g cm--2 sec-1)

773 3.0 x 10-1° 2.0 X io-io

873 1.2 x 10-9 1.2 X 10-9 1 VO

973 3.8 x 10-9 3.8 X M o

1073 - 1.5 X 10-5

1173 - 3.5 X 10"5

2. Fe - 10% Al

Temperature kp kl 2 (K) (g cm-4 sec-i) (g cm~2 sec-l)

1173 - 1.3 x 10"5

1273 - too rapid

3. Fe - 20% Al

Temperature kp kl 2 (K) (g cm~4 sec -1-) (g cm-2 sec-1)

1173 0 0 00 I o 1 — X

1273 Stage 1 4.1 1 00 I o — X

Stage 2 6.1 1

Stage 3 8.1 x 10-8 Pure Fe

------~ Temperature k P 2 -4 -1, (K) (g cm sec )

■ ■ -M ..■■■■ - - - - - 773 4.5 x 10"8

873 1.7 x 10"7

973 5.2 x 10"7 37

parabolic rate constants increased as the exposure time was increased. Kinetic data is summarised in Table 1.

(51) Nishida reported different kinetics from Strafford.

He studied the corrosion behaviour of Fe -3, 6 and 11 wt.%

Al in 1 atmosphere of S2(g) at temperatures between 823 K to 1173K. All the alloys corroded according to parabolic kinetics except Fe -11 w/o where an island growth mechanism was suggested as it was for Fe -10 Al by (45 46) Strafford ' . The corrosion rate was found to decrease with increasing aluminium content.

Although Nishida reported different kinetics from Strafford, it must be remembered that Strafford's experimental observations were made at an earlier stage of reaction than those of Nishida and their differing kinetics may be reconcilable. 5, 1»

F*S EARLY PARABOLIC ALLOY I STAGE

ALLOY LINEAR STAGE

2FeS — 2Fe’« S,» 4# (dissociation)

S,— 2S 1 (interface 2S . 4e — 2S* [ reaction*)

3S’» 2aT*— Al} S,

FeS

F.S-A.,5,

LINEAR STAGE nr ALLOY

Figure 6 Reaction Mechanism (schematic) 39

4.1.2 Description of Sulfide Scales and Their Formation Mechanism

(45 46) Strafford and Manifold ' found that the single layer of scales formed during the initial parabolic period of sulfidation of Fe -5% Al was very thin and easily damaged during the metallographic examination which made microscopic examination difficult. However the thicker scales formed under linear kinetics were duplex in nature consisting of an outer porous layer of Fei-yS containing a small amount of Al (~1%) in solid solution which was close to the outer scale - inner scale interface. The micro- porous inner layer of scale, which did not contain the striations or layers characteristic of the inner scale of iron, consisted of Fe1-yS mixed with finely dispersed Al2S3. The inner and outer layers of scale at temperatures between 773 K to 1073 K grew according to linear kinetics. At 1173 K, a non linear accelerating growth rate was observed for the inner layer causing it to be thicker than the outer layer.

The probable mechanism of sulfidation of the Fe -5% Al alloy is suggested (45 'to46) be similar to the theories regarding the behaviour of pure iron mentioned previously, but with some modifications. The mechanism is shown schematically in Figure 6 and will be summarised below.

The initial parabolic period of sulfidation suggests that the compact layer of FeS is growing by an outward diffusion 2 + of Fe ions towards the scale - sulfur interface 40

accompanied by an inward diffusion of vacancies in the opposite direction which are injected into the Fe alloy substrate. Thermodynamically, Al2S3 is also likely to be formed, but because of the fast diffusional growth of FeS relative to Al2S3, it will be buried under the FeS scales.

It has been reported 1 ' that FeS grows 10,000 times faster than Al2S3. The reason why Fe -5% Al sulfidises (4546) , slower than Fe is suggested ' to be due to Al3+ creating localised strain fields in the FeS lattice and acting as sinks for the Fe vacancies so reducing the vacancy diffusional rate.

Voids are formed at the sulfide - alloy interface in a similar manner to that described previously for pure iron except that here the small quantity of Al in the scale will strengthen the scale to make it more rigid and more difficult for plastic deformation to occur and so assists in the formation of voids.

The continued growth of the inner layer of FeS and Al2S3 is then possible by anistropic dissociation of the FeS around the voids to produce sulfur vapor (see figure 6 Stage II a).

Ingress of gaseous sulfur through grain boundaries and microfissures in the FeS scale helps to maintain a high sulfur potential within the FeS close to the scale - alloy interface to promote dissociation and to supply the sulfur to the voids for inner scale growth.

# Further growth of the outer Fei_yS scale then depends on the rate of Fe2+ ion diffusion from the alloy through the 41

inner scale which consists of Fe1-yS and Al2S3. it is suggested that the rate of growth of the outer FeS scale becomes linear because the rate of supply of Fe2+ ions is not fast enough to give the appropriate concentration at the inner scale - outer scale interface (Stage III, inter­ face 2) to achieve diffusional (parabolic) scale growth.

This is caused by the partial blocking effect of the dispersed Al2S3 particles in the inner layer which reduces the effective surface area for Fe2+ transport. It is suggested that eventually a steady state condition will be reached in which the rate of incorporation of the vacancies in the alloy and hence the rate of nucleation and growth of

FeS at the alloy - inner scale interface (Stage III, inter­ face 3) will control the linear growth of the inner and outer layers of scale.

At 1173 K, Fe -10% Al initially produced islands of sulfide which simultaneously expanded laterally and thickened resulting in a period of accelerating reaction rate. A duplex scale is ultimately formed. The outer layer consisted of Fe^.yS which contained a small amount of aluminium (4%) close to the inner layer - outer layer interface. The porous inner layer consisted of a fine dispersion of Fe S and Al,S, having the overall 1 — y composition of the double sulfide (spinel) FeS Al2s3.

The inner scale layers on the Fe -10% Al alloy contained more aluminium than those on Fe -5% Al alloy. At 1173 K and 1273 K, the thickness of the inner Al2S3 containing layer relative to the outer layer was more marked with 42

the 10% and 20% alloys than with the 5% Al alloys.

At 1273 K, Fe -20% Al formed a thin "whiskerlike" outer layer of Fe^ and internal sulfidation occurred forming needle shaped precipitates of composition approximating to

Al2S3. The needles penetrating into the interior of the alloy were orientated roughly normal to the alloy surface.

The mechanism of sulfidation of this alloy was not described but this process could be due to the phenomenon of discontinuous precipitation which has been the explanation for a similar morphology in other sulfidising alloys (49,50)

The linear kinetics and inferior resistance of Fe-Al alloys with Al contents greater than 5% at 1173 K and 1273 K were accounted for (45) by suggesting that due to the large molecular volume of Al2S3, the scale becomes more porous and less protective as the percent Al2S3 increases.

Another reason for the inferior resistance at the high temperatures is thought to be the low melting point of

Al2S3 (1373 K). Even an exposure temperature of 1073 K

( ~ 0.78 Tm) is perhaps high enough to allow rapid ionic diffusion in the inner layer and so the potentially protective effect of Al2S3 - containing scale is lost.

Thus this mechanism is quite different from that put forward for Fe -5% Al alloy at lower temperatures where the protective effect of aluminium was suggested to be due to

Al-S, formation which, since it is more stable than Fe, S, z J 1 -y accumulates in the inner layer and produces a blocking effect on diffusion of the reactant. 43

The island growth mechanism which was suggested for Fe -10% Al by Strafford was also reported by Nishida for Fe -11%

Al.

(51) Nishida also observed duplex scales in which the inner layer contained an accumulation of aluminium. As well as Fe1_ S and Al2S3, another compound which was similar to the

FeAl2S4 compound reported by Flahaut (53) was identified in this layer.

Nishida's explanation for the protective effect of aluminium was the same as that suggested by Strafford and Manifold (45,45) 44

!------. » /' t »/

!

• I// .

5 19 #f

9 fa i ' ii T*/ 100 *C r 517 • #- 4 a • J 516 M0*C ' 1000 *c 1 »09 . it: 0 Q2 OA Qi 06 10 M\4 mmrn tract in

Figure 7. Manganese concentration vs

^ lattice parameter of -manganous sulfide in 1 atm of sulfur vapor at (•) 1073 K;

(▼) 1173 K; (o) 1273 K; (^) 1373 K. (d4)

1 5.10

0 02 0A 06 Q0 10 man 6*cWi

Figure 8. Change in the lattice parameters with sulfur pressure (o) 1 atm; (A) 10-1 atm; (A) 10 4 atm; (•) 3.5 x 10~7 atm; all at 1273 (54) ««• i !T--^ -iiik’i 588?' -- i 5.80" -»r* , / T7 *>• 573^

I 5 76h v'1- <011

: ■ M 01 5 37?-- *-- * 34&-- --- <«*o 7J ^•TFr: TT“< ** * > *=\Vol 144*- a» 02 mot* fraction

Figure 9. Lattice parameters (a ,c ) of ferrous sulfide at various pressures. (o) 1 atm: (A) 10” atm; (54) (▼) 10 4 atm; (•) 3.5 x 10 ‘ atm; all at 1273 K.

I 600

1400

20 40 60 60 IOO MnS MnS

Figure 10. Equilibrium phase diagram of an FeS—MnS system 46

01 0 2 0J Q4 05 Ofi 0.7 08 09 10 Mn wit. n «ur1 mg alloy*

Figure 11. Changes in (Ns/NFe + N ) i-n sulfides with manganese content at various sulfur pressures at 1273 K (o) in 1 atm; (A) in 10 1 atm; (▼■) in 10 4 atm; — 7 (•) in 3.5 x 10 atm. The broken line (O) represents results in 1 atm of sulfur pressure at 1173 K (©) marks are from Burgmann et.al. (54) 4.2 Sulfidation of Fe - Mn Alloys

4.2.1 The FeS - MnS - S System

The lattice parameters of <*-- [Mn(Fe)]S formed at varying temperatures at a constant sulfur pressure of 1 atm are (54) shown as a function of manganese content in figure 7

It is seen that as the temperature increases, the lattice parameter of the <* - [Mn(Fe)]S decreases. The lattice parameter has a constant value for a particular composition. This value tends to move to a higher concentration of ferrous sulfide as the temperature increases and the pressure decreases. In figures 7 and 8, the two parallel dashed lines separate the existence of the o'- [Mn(Fe)]S single phase region from the two phase region of

oc - [Mn(Fe)]S and [Fe(Mn)]S. The composition at which the lattice parameter has a sharp break shows the solid solubility of ferrous sulfide in oc - [Mn(Fe)]S.

The pseudobinary diagram of the FeS - MnS system in figure 10 shows the solid solubilities of MnS in ferrous sulfide and FeS in - [Mn(Fe)]S. It can be seen that the solubility of FeS in MnS depends strongly on the sulfur pressure. Results by Nishida (54) (shown by the broken lines), revealed that decreasing the sulfur pressure from

1 atm to 10 atm increases the solubility of FeS in MnS by about 25 ~ 30 mole % MnS.

Figure 11 shows the mole ratio, sulfur to metal [N /(N + N )] in sulfides equilibrated with Fe - Mn s Fe Mn alloys. The single-phase regions of solid sulfide solution 48

based on ferrous sulfide and on oc - manganous sulfide are indicated by the shaded areas on the diagram. It can be

seen that the mole ratio is always higher than one in both

single phase regions, and that in the ferrous sulfide solid solution region, it possesses a constant nonstoichimetric composition depending on the sulfur pressure. However, in the two phase region, the ratio decreases as the ~ manganous sulfide content decreases and it is almost equal to one in the oc-manganous sulfide solid solution. This, of course, is a consequence of the small degree of nonstoichimetry found in oc- Mn.^ S. Mixtures of the two phases posess sulfur/metal ratios which reflect the properties of the metals in the mixture. 49

4.2.2 Kinetics

Binary Fe - Mn alloys ranging from 0 - 100 w/o Mn (where w/o means weight percent of) were exposed to 1 atmosphere of pure sulfur vapor at temperatures of 973 to 1273 K (54 '

55) The sulfidation rate and scale morphology depended on the alloy composition as seen in Table 2.

Alloys containing up to 11 w/o Mn reacted at a rate independent of alloy composition and at the same parabolic rate as pure iron.

As the manganese concentration was increased, the parabolic rates of corrosion decreased sharply, a trend which continued up to 63 w/o Mn. The activation energy increased continuously from a value similar to that for pure iron sulfidation to a value similar to that for the sulfidation of pure Mn(135 kJ mole 1).

For alloy levels of between 63 - 80 w/o Mn, the corrosion rate did not closely follow the parabolic law. For alloys containing above 80 w/o Mn, the parabolic sulfidation rate was again independent of the composition of the alloy in this region and was equal to that of pure Mn.

The sulfidation of pure manganese has also been studied in _ s — 4 H2/H2S atmospheres, corresponding to 10 ^ P ^10 atm

/ r r \ b^2 at temperatures between 1073 and 1273 K . The reaction kinetics were found to be parabolic at all sulfur pressures (see Table 3). The parabolic rate constant was proportional to P where n = 6.1. Reaction rates were two to three b2 50

orders of magnitude slower than for pure iron under the same conditions. 51

Table 2: Rate Constants obtained during sulfidation of

Fe - Mn alloys, Pure Iron and Pure Manganese

Mn content 0-11 w/o

Temperature k p . 2 -4 -1, (K) (g cm sec )

973 2.0 x 10-6

1073 5.0 x 10-6

1173 1.0 x io“5

Fe - 29 w/o Mn Alloy

Temperature k p -1, (K) (i g 2 cm"4 sec ) 1 * -vj — 973 4.0 x i o

1073 1.0 x IQ'6

1173 3.0 x 10"6

Fe - 48 w/o Mn Alloy

Temperature k p -lx (K) (g cm sec ) CO 1 ! o - 1 973 7,9 x

1073 3.2 x 10 7

1173 1.0 x io~6 & i 1 o —

1273 2.0 x 1 Table 2

Fe - 58 w/o Mn Alloy

Temperature k P (K) (g( 2 cm "4 sec -Is) CO 1 I o —

973 3.2 x 1 c- 1 1 o — 1 1073 1.3 x 1 1 I- 1173 4.0 x o 1 rH 1273 1.0 x o

Fe - 75 w/o Mn Alloy

Temperature k P -1, (K) (gi 2 cm~4 sec )

973 2.5 x 10"9 -9 1073 6.3 x 10 00 h-* 1173 4.0 x o r- i i o —

1273 1.0 x i

Fe - 8 0 w/o Mn - Pure Mn

Temperature , 2 -4 -14 (K) (g cm sec )

973 7.9 x 10"10

1073 3.2 x 10"9

1173 1.6 x 10-8

1273 1.6 x 10”8 53

Table 3; Values of the Parabolic Rate Constant,

2 -4 k (g Scm sec ) , for Growth of °<--Mns Pp (56) on Mn

Temperature K P , atm b2 1073 1173 1273

-8 -10 -9 -9 10 7.6 x 10 1.9 x 10 3.8 x 10 - 7 -9 -9 -9 10 1.2 x 10 2.9 x 10 5.5 x 10

2.8 x 10"9 -6 -9 -9 -9 10 1.9 x 10 4.3 x 10 8.1 x 10

8.0 x 10 -5 -9 -9 -8 : 10 2.2 x 10 5.1 x 10 1.2x10 -8 5.2 x 10 9 1.2 x 10 -4 -9 -9 -8 10 3.5 x 10 9.0 x 10 1.9 x 10

3.5 x 10"9 54

4.2.3 Description of Sulfide Scales and their Formation Mechanism

Alloys containing up to 11 w/o Mn produced single phase scales consisting mainly of pure Fe S and some precipitates of Mn(Fe)S (54) . Manganese was concentrated mainly toward the scale - gas interface. However, as shown in the

FeS-MnS phase diagram in figure 10, the solubility of manganese sulfide in ferrous sulfide is 4 - 7 mole % MnS, depending on the temperature. The manganous sulfide content in the scale is within the solubility limit for ferrous sulfide for 1 w/o Mn alloy. According to Wagner (57) , the different distribution of alloy constituents in the scale is due to the difference between the self diffusivities of each metal ion. When the ratio of self diffusivities of iron and manganese (D /D ) was assumed ^ Fe Mn to be 0.42, the observed concentration profile was ( 54) consistent with the calculated one . Thus the self diffusivity of manganese ions is 2.4 times higher than that of ferrous ions in this scale.

The manganese content in the scale increases with the manganese content in the alloy. When the Mn content near the scale - gas interface is greater than the solubility of MnS in the FeS phase, an oc- manganous sulfide phase starts to precipitate. The initial precipitate is concentrated near the scale surface but with increasing manganese concentration in the alloy, the precipitates spread through the scale. These precipitates were leaf­ like in appearance and parallel to the scale surface. 55

Nishida (55) reported that the Mn content dissolved in the

FeS scale reached a saturation limit for a 6 w/o Mn alloy.

From these results, it can be seen that the kinetic factor

(self diffusivity of each metallic ion in the scale) has a great effect on the site of precipitation in the scale. Since the scales on the alloys containing up to 11 w/o Mn contained very little MnS, the corrosion rates were very similar to that of pure iron, the rate being controlled by (55) diffusion in the FeS phase

Scales formed on alloys with 11 to 63 w/o Mn were duplex, with the more stable cubic Mn(Fe)S under an outer layer of hexagonal Fe S. The scale was thick for alloys with l-y greater than 30 w/o Mn and the rate of growth of each layer was reduced with increasing manganese content. As the temperature increased, the relative thickness of the inner layer also increased, e.g. at 1273 K, Fe -58 w/o Mn alloy had no outer layer, only an inner layer being formed.

Results of Arrhenius plots of growth - rate constants for inner, outer and total layers indicates that the activation energy for the rate constant of the inner layer is greater than that of the outer layer (54) . The scale on Fe -58 Mn consisted largely of single phase [Mn(Fe)]S but contained ferrous sulfide precipitates near the scale surface. This is due to the low sulfur activity near the alloy - scale interface, and consequently high solubility of FeS in MnS. However, as the sulfur activity rises toward the scale - gas interface, the solubility of FeS 56

is decreased and phase separation occurs, as predicted

from figure 10.

The scales formed on the alloys containing between 63 and

80 w/o Mn will now be discussed. Sulfidation of Fe - 75

w/o Mn alloy results in a scale consisting of single phase

oc - manganese sulfide with some dissolved iron. The ferrous sulfide phase dissolves into the scale, the iron

concentrating significantly toward the scale - gas interface. This distribution is the opposite of that found in scales on low Mn alloys. This, according to the Wagner theory suggests that DFe is greater than D in the

manganous sulfide, but that there is not yet enough evidence to verify this data ( 54)

The alloy substrate consists of mixed phases of cK and with only the phase close to the surface. The scale on an Fe - 80 w/o Mn alloy contains hardly any iron. As sulfidation of Fe - 75 w/o Mn alloy continues, the scale consists of a single oC- manganous sulfide phase similar to 80 w/o Mn, while iron atoms diffuse into the alloy substrate through the pphase in the alloy. However, when

the concentration of iron in the ^3 phase increases, the JB phase transforms into the phase and the corrosion subsequently is similar to that of the Fe - 58 w/o Mn alloy. Not surprisingly, the corrosion of this alloy does not follow parabolic kinetics.

The corrosion rate of the 80 w/o Mn alloy was found to be equal to that of pure Mn. Since the corrosion rate of 57

(54) 80 w/o Mn is so low, the diffusion process in the

alloy substrate as well as the scale diffusion process has

to be considered. The substrate initially consisted of the

jg> - Mn phase in which alloy diffusion rates are thought to

(54) be high . The MnS scale formed contained very little

iron. This is because the sulfidation of manganese is O thermodynamically more favourable since AGf for MnS is

O than AGf for FeS as shown in the values below:

0 Mn + *2 S2 = MnS AGf = -67,200 + 16.17 T cal/mole S2

O Fe + h s2 = FeS AGf = -39,198 + 14.6 T cal/mole S2

The sulfidation of pure manganese resulted in the

formation of a single phase scale of cubic OC- MnS at all

sulfur pressures. The sulfidation results are consistent with the fact that cx. - MnS is a metal - deficit semi­ conductor. In oC- MnS, the point defects of interest are cation vacancies and positive holes, therefore the

following solid-gas equilibrium is considered ;

Y " S + VM + 2, h S, (g) s Mn ...(24)

(60) According to Wagner , the parabolic rational rate constant k for the growth of the OC - MnS scale can be described with equation (25) which is derived from equation (11) mentioned previously.

eq \ '2 D_. d In P_ 5: Mn S .••(25)

Where = self-diffusion coefficient of manganese in

(X- MnS and P* is the sulfur pressure at the metal - z 58

scale interface corresponding to the dissociation pressure of ^ - MnS (61), p" is the sulfur pressure at scale - gas interface and C is the average concentration in eq equivalents of manganese in oc- MnS.

Now using the relation according to equation (24) for the manganese self diffusivity

1/6 (26) Mn Mn where D is the self-diffusivity at P_ 1 atm. Now Mn J S substituting equation (26) into equation (25) the following equation is obtained: P" C D S eq Mn 1/6 d In

ps 1/6 C D eq Mn d PS

ps.

ps C D p -5/6 p eq Mn S2 d S2 . . . (27)

ps 59

Now integrating equation (27) gives:

P" S 1/6 C D eq Mn

ps.

-PS 1/6 6 C D eg Mn

ps

P " 1/6 pi 1/6 3 C D I S0 — S0 eq Mn V 2 2

» 1/6 3 C D 1 — V/s ...(28) eq Mn pSj

The rational rate constant k (in equivalent cm 1 s 1) is

2 _ A _ 2. related to the parabolic rate constant, (in g cm s ) by the following equation:

h V k / (M ) eq p s .•(29) where is the equivalent volume of MnS, k^ is the parabolic rate constant and is the atomic weight of S

2 k (M ) r s . . . (30) 60

Substituting equation (28) into equation (30) gives

O 2 /P 1 / 6 6 C D (M ) _ 1/6" / Si , _ ___eg Mn s P c / -| _/__ 2 a ...(31) P V S2 U PCJ '

Now dimensional analysis of the terms C and V are J eq eq shown below:

Mass of o<.-MnS _ pMnS MA(I x Volume M.. Mns Mns

MnS and where is the molecular weight Mns r e<2 ?Mns of

Therefore

2 0 /M 0 MnS ...(32) M/0 2 \f[ MnS

Now substituting equation (32) into equation (31) gives:

l/6> 6 D P 1/6 Ms 9Mns (l-(PS2/PS2)

cm sec Mn S2 “~M------Mns

... (33)

Since (P' /P" )

Figure 12

T(K)

0* 21,000 tt 2000) cat /mot*

Q*3I,000(± 3,000)\ cal/moU

0 ••7,200 (±500) ca(/mal« A NISHIDA el a| O OHTA «t 04

Arrhenius plots of the self diffusion coefficient of Mn in the oc-MnS scales formed in flowing H2S-H2 at 1 atm pressure , static H2S-H2 atmospheres at 0.1 atm pressure, and in static pure sulfur vapor at 1 atm pressure. 62

experimental results.

In an attempt to compare their own findings with previous research on the sulfidation of manganese which was carried out in one atmosphere of pure sulfur vapor ^4) an^ a static H2S-H2 atmosphere at a total pressure of 0.1 atmosphere , Elrefaie and Smeltzer ^6) used the parabolic rate constants at different sulfur pressures to o work out D using equation (33) and the thermochemical Mn data for oc -MnS and H2S-H2 equilibria ^6). The

O variation of these values of D with temperature is shown Mn in figure 12. The diffusivity values agree only within a factor of five. One cannot say whether the differences among the results come from the different exposure conditions and/or the quality of the manganese metal. Elrefaie and Smeltzer suggest that the manganese self diffusion coefficient represents an average value for the metal lattice diffusion by point lattice defects, and by grain boundary and dislocation paths in the polycrystalline oc-MnS scales. 63

(6) Table 4; Sulfidation Rate Constants of Fe-25Mn-5Al Alloy

2 4 kp(g S/cm sec)

S2 700°C 8 0 0 °C 900°C (atm) 1 CO h-> o 3.6 x 10"10 4 x 10_1° 9.9 x 10_1° CO 1 1 o —

1 1.0 x 10“9 1.1 x 10"9 1.7 x 10"9

2 k1(gS/cm sec)

10 7 5.5 x 10~6 2.2 x 10"6 1.6 x 10“5 1 1 o —

1 1.5 x 10"5 1.9 x 10”5 1.2 x 10“5

10 5 2.2 x 10"5 3.9 x 10-5 3.9 x 10~5

io"4 2.9 x 10"5 8.5 x 10"5

/ /T O \ Table 5: Parabolic Rate Constants of Fe-25Mn-20Ar

2-4-1 k (kg m s j P S2 (Pa) 973 K 1073 K

102 9 x 10-9

10 9 x 10-9 sigmoidal

1 sigmoidal l — 1 i o — 1

i 5 x 10“9 irreproducible CO 1 CO rH o 10-2 X

10“3 2 x 10 8

10“4 1 x 10"8 64

5,0 THE SULFIDATION OF Fe-Mn-Al ALLOYS

There are very few reports available on the sulfidation of iron based alloys containing both manganese and aluminium.

However, the oxidation of such alloys has been studied by a number of experimenters (62,63,64).

Tomas K reported some promising information on an Fe-30Mn-8Al alloy. At a temperature of 978 K and — 9 P = 3 x 10 atm, the alloy was found to be more corrosion b2 resistent than the stainless steel types 304, 316 and 310. The scale formed consisted mainly of MnS overlying a fine layer of aluminium rich sulfide thought to be either Al2S3 or

FS A12S4-

f a r O \ Sulfidation studies ' ' on Fe-25 Mn-Al alloys containing 5 and 20 wt % Al have been carried out at 973 K, 1073 K and 1173 K at pressures between 10 atm and 10 atm.

5.1 Kinetics

The sulfidation kinetics of Fe-2Mn-5Al at the three — 9 — Q temperatures investigated and P =10 and 10 atm were b2 found to be parabolic for the investigated periods ranging up to 360 minutes. At 973 K and P = 10“ 7 atm, the 2 kinetics were parabolic up to 160 minutes, then trans­ formed to rapid linear behaviour. These kinetics were also found at higher sulfur pressures up to 10 4 atm. With increasing sulfur pressure, the kinetics became linear 65

earlier. These type of reaction kinetics were found for

Fe-25Mn.-5Al at 1073 K and 1173 K, 10_6

and linear rate constants were calculated by approximating

the kinetics to these relationships as defined by the shapes

of the curves obtained.

The sulfidation kinetics of Fe-25Mn-20Al alloy

measured at 973 K and P = 10 3, 10 4 and 10 6 atm were b2 found to be parabolic after an initial period of 16.7 hours

which decreased as the P was increased. At 1073 K and

P

short2 initial periods of 3 minutes. At P = 10 -6 atm, b2 very irreproducible kinetics were observed, both parabolic

kinetics and faster, less protective kinetics being reported.

At 973 K and 1073 K, the parabolic rate constants were

relatively independent of P , the magnitudes tending to b2 decrease as P decreased as shown in Table 5. b2

An Fe-25Mn-20Al alloy was found to sulfidise according to

sigmoidal kinetics at 1073 K and P = 10 5 and 10 4 atm b2 . After an initial period the rate accelerates rapidly

and then slows down to a relatively slow rate. Rapid

kinetics were found to occur at an earlier period at the

higher P . Onset of the final period of slow sulfidation b2 also occurs earlier at the higher P . When the period of b2 rapid sulfidation was approximated by a linear rate

equation at P =10 4 and 10 5atm, k-. was found to be 10 4 O g -L and 3 x 10 - 5 kg m- 2 s-1 respectively. 66

Smeltzer observed that of the alloys investigated at

1073 K, Fe-25Mn-20Al showed the best corrosion resistance but this was limited to P 10 6 atm, the kinetics being b2 X sigmoidal at higher sulfur pressures. This alloy was even more corrosion resistant than the binary Fe-28Mn and Fe-50Mn alloys . 5.2 Description of Sulfide Scales and Their Formation Mechanism

_ g At low sulfur pressures, P

- MnS containing precipitates of MnAl2S4« A small amount of internal sulfidation was observed beneath the scale. The extent of internal sulfidation was found to

increase with increasing temperature. The internal sulfides were composed of MnAl2S4 with Al2S3 as a secondary minor phase. Since Fe-25Mn-20Al alloy is more corrosion resistant than the binary Fe-28Mn and Fe-50Mn alloys at 1073 K, Smeltzer ^ suggests that aluminium in the ternary alloy does have a useful sulfidation effect probably caused by a diffusion blocking effect of MnAl2S4 precipitate in the MnS scales and internal sulfidation

zones.

The Fe-25Mn-20Al specimens which sulfidised parabolically produced a scale containing only oc- MnS. The scale consisted of a thin compact protective inner layer and a dispersed outer layer. A continuous product layer of

oc- MnS would be expected since oc- MnS is relatively more

stable than Al2S3 and FeS under these conditions .

At 1073 K and high sulfur pressure, P 10 7 atm, 2 sigmoidal kinetics were observed due to the failure of the initially formed protective scale leading to the formation of large nodules beneath the thin c<. - Mn(Al)S layer. The nodules then spread to form a relatively thick uniform 68

scale which consisted of the hexagonal mixed sulfide

M Al2S4 where M is Fe and Mn. After a lengthy reaction period, a trace of Al2S3 was detected. When a nodule was initially formed, it had a homogeneous and fine grained appearance. However, as reaction time was increased, the nodules tended to have a fibrous structure with the fibres orientated normal to the alloy surface. After 24 hours — 4 reaction at 1073 K and P = 10 atm, small needle-like 2 protrusions of sulfide into the alloy surface were observed. The nodules were composed mainly of (Fe,Mn)Al2S4 and

©c - Mn(Fe)S with the innermost sulfide layer consisting of a mixture of Al2S3 and (Fe,Mn)Al2S4. Electron probe

Microanalysis results showed that the alloy subsurface region was greatly depleted of both Mn and Al.

Alloy depletion in Mn could lead to the formation of Al2S3 at the scale base if MnS was destabilised with respect to

Al2S3 via the reaction:

3 MnS + 2 Al = A12S3 + 3 Mn

This reaction is thermodynamically favoured at T = 1073 K when the activity ratio (a A/I )3/ (a ,) 2

An inert platinum marker was found on the outer Mn(Al)S layer of scale when the alloy was sulfidised at 1073 K and

— 4 P = 10 atm for 98 minutes. This indicates that the b2 outer layer transmitted sulfur. It was concluded that the growing nodule pushed the preformed Mn(Al)S layer and the Pt marker outwards. Thus the position of the marker 69

gave no information on the transport mechanism through the

nodules and the fully developed scale. When a marker was

placed between a Mn-34Al alloy and oc- MnS diffusion

couple and annealed in an evacuated silica ampoule for 74

hours, the marker was found in the product layer under a

thick MnAl2S4 layer and separated from the alloy by a thin

duplex layer of MnAl2S^ and Al2S3. This indicates that the

layer of MnAl2S4 and MnS grows by metal transport and it would similarly be expected for the MAl2S4 and Mn(Fe)S

scale layer.

It has been suggested that failure of the initially

formed protective Mn(Al)S duplex scale containing Al2S3,

leads to nucleation and growth of the nodules. The overall rate of reaction accelerates as the nodules expand

laterally beneath the thin scale. As the nodules impinge on one another their surface coverage approaches unity and the sulfidation rate becomes similar to that for the growth of a uniformly thick sulfide scale. Diffusional growth of the fully developed scales was possibly largely determined by the transport properties of the innermost Al2S3 and

(Fe,Mn)Al2S3 layer.

The usual way of treating nodule growth kinetics is by deriving the surface area fraction covered by sulfide, 0, using statistical methods (69,70) . The result obtained is

0=1- exp (- K Pm tn) ... (34)

where K is a constant, P is the reactant pressure, and m and n are positive exponents whose values are determined (70 71 72) by the detailed mechanism of nodule spread '

Calculation of the kinetics of weight uptake may be obtained from this model if the degree of nodule impingement is low and the nodule shape known. If a nodule has a circular base and spreads laterally at a constant rate, its base radius, R is given by:

dR/dt = V ...(35) where V is a constant. If the rate of thickening of a nodule is assumed to be diffusion controlled, then its thickness at any point, X, is given in terms of the parabolic growth rate by:

X = k (t — t ) ...(36) P c where t is the time at which the nodule periphery reached the point in question and sulfide growth commenced. There fore for a nodule which nucleated at t = 0

0 < t < Rt , t = r/v ...(37)

From equations (36) and (37) :

X = k (t - r/v) ... (38) which describes the profile of the nodule. Volume integration then yields for the nodule volume: 71

V = TT R 5/2 (^)h •••<39) 15 v V

If the area density of nodules is N, and if the degree of impingement is small, then

9 = NTT R2, AW /A = N £ V ...(4 0)

whence

k 8 e5/4 .••(41) 15 N v ) where AW/A is the weight uptake per unit area and is the sulfide density. Substituting equation (34) into equation

(41) gives

AW „ ri . m.n.,5/4 n c ,.,^,3/2 1/2 = 8 ^[1 - exp (-K p t )] 15 (NTT) v

.. . (42)

which describes a sigmoidal rate. Double differentiation of equation (42) enables t^, the time at which the inflexion in the rate curve occurs, to be defined. It is found that:

ti °° ps~m/n • • *(43)

which is in agreement with the observed increase in t^

(from 92 to 182 minutes) with decrease in P (from 10 to b2 1 Pa) at T = 1073 K. This simple model breaks down at high coverages and cannot then be used to analyse the observed

73

Table 6: Comparison of Sulfidation Rate Constants^54'55'88)

973 K

2 4 kp ( g /cm sec ) PS2

(atm) Fe-25 Mn-20Al Fe-25 Mn-5Al - - Fe-50 Mn Fe-28 Mn i rH o 9 x 10'11 4.5 x 10“9 9.8 x 10"8

10 5 io‘6 5 x 10"11 " 1.5 x 10~9 6.4 x 10~8 Table 7: Comparison of Sulfidation Rate Constants (54,55,68)

1073 K

2 4 Ps2 fcp ( g. /cm sec )

(atm) Fe—25 Mn-20A1 Fe-25 Mn-5Al Fe-50 Mn Fe—28 Mn

10~7 3 x to'10 1.2 x 10'9 3 x IQ'8 1 00

O 2 x tO”10 . .1.1 x 10'9 1.2 x 10'9 9.7 x 10'9

10~9 1 x tO"10 4 x ICT10 1.2 x IQ'9 75

5.3 Comparison of Sulfidation Rates

Due to the complicated sulfidation behaviour of Fe-Mn-Al alloys, only the alloys undergoing parabolic sulfidation will be compared. As shown in Table 6, Fe-25 Mn-20A1 displayed the highest sulfur corrosion resistance compared with Fe-50Mn at T = 973 K and P = 10 6 atm and -4 2 P = 10 atm. The Fe-50Mn alloy was found to sulfidise z at about one order of magnitude less than Fe-25Mn alloy and at least two orders of magnitude less than the sulfidation rate of pure iron.

From Table 7, it can be seen that the sulfidation rates of Fe-25Mn-20Al alloy at T = 1073 K at P <

PART B - EXPERIMENTAL WORK

6.0 EXPERIMENTAL PROCEDURE

6.1 Introduction

Experiments were carried out on Fe-4.5Mn-8.8A1-0.36C alloy (composition in wt %) at temperatures in the range 973 K to 1173 K in H2S/H2 atmospheres corresponding to

10 8 atm ^ P

6.2 Experimental Conditions

The minimum sulfur partial pressure (P ) required to form b2 each sulfide (FeS, FeS2, MnS and Al2S3) at 973 K, 1073 K and 1173 K was calculated using thermodynamic data from the literature (73) . An error results from the fact that sulfide intersolubility was not taken into account. In each case the reaction is of the form

xM(s) + S 2(g) = MxS2 (s) and for unit activity of the solid phases

1 = K = exp ( — AG°/RT) ...(44) P P b2

The sulfur partial pressures were controlled by mixing

H2 and H2S in fixed proportions, the required H2/H2S ratio being obtained from the known standard free energy change 77

(73) Table 8: Standard Free Energies of Sulfide Formation

o Reaction A Gf Temperai K (J/mole S,)

H2(g) + j S2(g) H2S(g) - 4.23 x 104 973 - 3.74 x 104 1073 - 3.26 x 104 1173

Fe(s) + ± s2(9) = FeS(s) - 9.97 x 104 973 - 9.47 x 104 1073 - 8.97 x 104 1173 CM i O r- 00 LD — + X FeS(s) + 2 s2(9) = FeS2 (s) 1 973 + 1.06 x 104 1073 + 1.96 x 104 1173

Mn(s) + -^ S2(g) = MnS(s) - 2.05 x 105 973 - 1.98 x 105 1073 - 1.92 x 105 1173

■| Al(s) + S 2 (g) = Al 2S 3 ( s) - 1.92 x 105 973 - 1.82 x 105 1073 - 1.74 x 105 1173 78

of the equilibrium:

H2S (g) = H2 (g) + \ S2 (g) ... (45)

Because of the excellent agreement found for iron sulfidation ^4) between experimental observations and

Wagner's theory, regardless of whether hydrogen is present or not, it was assumed that the presence of hydrogen had no effect on the sulfidation reaction.

The free energies for each reaction are shown in Table 8. 79

The equilibrium values of P for each sulfide are shown in b2 Table 9 together with the corresponding H2S/H2 ratios.

Table 9: Equilibrium Data for Sulfide Formation

PS Temperature Substance b2 Formed ph2s/ph2 K (atm)

973 FeS 2 3.81 x 1011 1.15 x 108 FeS 2.00 x 10"11 8.41 x 10."4 Al2S3 2.60 x 10'21 9.61 x 10~9 MnS 1.13 x 10"22 2.01 x 10"10

1073 FeS 2 5.50 x 1011 4.92 x 107 FeS 6.10 x 10~10 1.64 x 10-3 A12S3 1.60 x 10"18 8.50 x 10"8 MnS 5.26 x 10"2° 1.52 x 10"8

1173 FeS 2 7.45 x 1011 2.44 x 107 FeS 1.89 x 10"9 1.23 x 10"3 A12S3 3.40 x 10"16 5.23 x 10"7 MnS 9.06 x 10"18 8.52 x 10“8 80

The chosen experimental conditions for reaction are shown in Table 10. Comparisons of Table 9 and 10 show that under all conditions, the sulfides FeS, Al2S3 and MnS were expected to form, but not FeS2.

Table 10; Experimental Conditions

Temperature h2s/h2 PS2 (K) (atm)

- 4 973 1 x 10 1.8700

1 x 10"5 0.5910 -6 1 x 10 0.1870

1 x 10"7 0.0590

1 x 10"8 0.0190

1073 1 x 10"4 0.6632

1 x 10"5 0.2097

1 x 10"6 0.0663

1 x 10"7 0.0210

1173 1 x 10"4 0.2829

1 x 10"5 0.1340

1 x 10"6 0.0280

1 x 10“7 0.0089 81

6.3 Sample Preparation

The experimental alloy, Fe-4.5Mn-8.8Al-0.36C was produced by the School of Metallurgy, University of N.S.W.

The alloy was received in rod form. After the specimen was sliced into discs on a diamond saw to approximately

0.4 mm thickness and 12 mm diameter, 2 mm diameter suspension holes were drilled into each disc. The sample faces were ground down successively to 600 grade abrasive paper, then polished with 6 urn diamond paste and then finally with 1 urn diamond paste until mirror like surfaces were obtained. They were then washed with water and soap, water, acetone and methanol. The samples were then stored in methanol until required. Prior to a test run, each sample was given a quick polish with 1 urn diamond paste on all faces to remove any oxide or dirt that may have formed or collected on the sample which might have hindered sulfidation. Its dimensions were measured and it was then cleaned in soap solution followed by methanol. When dry, the sample was weighed, placed vertically on a small platinum hook and then immediately hung on a fine silica thread inside the reaction tube. Figure 13. Schematic Assembly of Sulfidation Rig Used

A Gas Regulators I Rotary Pump B Drying Column containing J Vacuum Gauge anhydrous magnesium

perchlorate K Gas Bubbler containing NaOH C Rotameters L Gas Outlet D Vycor Glass Reaction

Vessel M Vent N Specimen E Ni-Span C Spring F Winch 0 Gas By-pass G Tube Furnace P Positioning Forks H Cathetometer 83

6.4 Sulfidation Apparatus

The apparatus used in these experiments is shown in figure 13. The specimen was suspended from a McBain balance with a fine silica thread. The weight change (AW) in the sample undergoing sulfidation was followed continuously by observing the change in the length of the spring on the McBain balance with a cathetometer.

The Ni-Span C spring, whose modulus of elasticity was almost independent of temperature, was calibrated by loading different weights on it. The distance it moved was measured with the cathetometer everytime the weight was added or removed. The calibration value was found to be 0.040 g/cm ± 2% extension. The precision of the cathetometer was 0.01 mm, yielding an overall balance sensitivity of 40jUg.

The temperature of the furnace was measured by placing a

Pt Vs Pt - 13% Rh thermocouple between the reactor and the furnace mullite tube. This thermocouple was connected to a P.I.D. controller. Temperature control to within ± 4°C was achieved.

Separate rotameters were used to measure the flowrates of

H2S and H2 gas which were controlled with needle valves.

Rotameters were calibrated using the soap bubble method.

A N2 flow was provided in order to prevent sulfurous gases entering the spring chamber (see figure 13). H2S and H2 gases were dried by passing through anhydrous magnesium 84

perchlorate towers before entering the reactor. The effluent gas from the bottom of the reactor tube was bubbled through 30% NaOH solution to remove all traces of hydrogen sulfide before venting.

The entire reactor was constructed from vycor glass and designed to be vacuum tight thus preventing any air contaimination during sulfidation. The Quickfit connections were all made air tight by lubricating with high vacuum grease.

After the prepared sample was loaded onto the balance, the system was evacuated with a vacuum pump then back filled with nitrogen. This procedure was repeated and then nitrogen gas was allowed to flow continuously.

The use of a gas by-pass network enabled isolation of the specimen within the reactor permitting the gas mixture to be adjusted and maintained prior to the commencement of each run. The power input to the furnace was adjusted until the required temperature was reached. On reaching this temperature, the flow of H2S/H2 was redirected from the by-pass network into the preheated reactor, the sample was lowered into the reactor tube by the use of a winch and two positioning forks. At this instant, the stopwatch was started and the zero reading of the spring was obtained with the cathetometer. The extension of the spring was assessed initially every couple of minutes and at longer time intervals in the later stages of reaction. At the completion of the run, the sample was winched up and 85

allowed to cool to room temperature while the system was purged with nitrogen gas. 86

6.5 Characterisation and Identification of Scales

After each alloy had been sulfidised, x-ray diffraction analysis of the sulfide scales was undertaken,followed by microscopic examination and, in some instances, electron probe microanalysis.

6.5.1 X-Ray Diffraction Analysis

X-ray diffractometry of the outer sulfide scale "in situ" was performed i.e. the x-ray beam was directed onto the sulfide scale whilst still in contact with the alloy. Cu - K^ radiation (A = 1.54178A) was used. The patterns were identified by comparison with ASTM standard diffraction card "d" spacings. These appear in the Appendix along with data from the reference diffraction used. Intensities were calculated but oxides have been shown to have a preferred crystallographic orientation (74) with respect to the surface on which they lie . This (75) results in a significant difference between observed intensities and those given on ASTM cards which represent patterns for specimens composed of randomly orientated crystals. This was taken into consideration when identifying the diffraction peaks obtained, as precise correlation with standards was rare.

6.5.2 Metallographical Examination

X-ray diffraction analysis complete, the corroded specimen was mounted on edge in cold setting resin, using vacuum impregnation. Conventional metallographic polishing 87

methods were applied using 400, 600 and 800 grade abrasive paper until a cross section was obtained. Specimens were then polished on nylon cloth impregnated with 6 jum diamond paste, followed by further final polishing with 1 ^um diamond paste. The scale - metal cross section was then examined under a microscope and the scale structure was photographed.

6.5.3 Electron - Probe Microanalysis

The distribution of elements within some scales was investigated by electron probe microanalysis. The results quoted have been corrected for paralysis time, background radiation, absorption, atomic number and fluorescence effects. The surface of each specimen was coated with carbon to ensure good electrical conduction. Standards used were pure metals and CuFeS2. The instrument used was an ARL Model EMX employing a takeoff angle of 52.5° and an electron accelerating voltage of 25 KV. 6.6 Analysis of Data

Parabolic sulfidation rate constants were estimated by regression of the appropriate data on the equation:

W' .. . (46)

2 _ 4 time t and is the parabolic rate constant (g cm sec

Linear rate constants were obtained by regression of data on the equation

^ + W' = k1t ...(47)

— 2 -1 where k^ is the linear rate constant (g cm sec ).

The Activation Energy, E , for the reaction was calculated using the Arrhenius equation:

-E_/RT k = A e A P where A is a constant, R is the gas constant and T is the absolute temperature. By plotting ln(k ) against the reciprocal of the absolute temperature, the slope of the plot gives E /R. CO

QD atm IX (•W CD 3 * 0 LD S/DW)

oo I r-H o 4-1

rti £ ro b/w

CN

V -=r O LD LD ND ID CD o O O O CN o ro O LD O O O O CN LD O o V 21 1 LU 2: 21 — — I i

•H g Figure 1 4: Linear plots, weight gain vs time, for the sulfidation of the alloy at 973 89

© Q

© ♦

oo l rH o Oi W

n atm ^W3

4DW

^tB/WVl ©

■€> un I o 4-1 cu s f0 CO O n S 1 4.

LD ■C" O o ■c* o LD o ro o ro o LD o CM O CM o LD o o o LD o TIME (MIN) Figure 15: Parabolic plots, mg2 cm 4 vs time (min) for the sulfidation of the alloy at 90 atm . '_W3

*DW

^(b/WV)

Figure 16; Parabolic plots mg2 cm vs time (min) , for the sulfidation of the alloy at 91 f v i 4 fd g -> CD O

LD O 2W0 O

NO O

2

CH/WV1 CN O O K) O O CN O CM O LD o LU x;

Figure 17: Parabolic plots, mg2 cm vs time (min) , for the sulfidation of the alloy at 92 93

7.00 RESULTS

7.1 KINETICS

Figure 14 shows plots of zlM/A versus time for specimens sulfidised at 973 K and 10 8 atm

-7 At the higher sulfur partial pressure of 10 atm the rate initially accelerated and then the reaction kinetics tended

2 to become protective. A plot of (Am/A) versus time for the specimen is shown in figure 16. Parabolic plots of the same data confirmed parabolic kinetics. The parabolic rate constant was 7.8 x 10 g cm s

At 973 K and P = 10 6 atm, the specimen reacted initially z at a very rapid rate followed by slow protective kinetics. Parabolic plots of the same data are shown in figure 17 and indicates an intermediate type behaviour.

The reaction kinetics at 973 K and P = 10 -5 atm were b2 initially protective upto 170 minutes after which the rate began to accelerate. These protective kinetics were not parabolic as shown in figure 15.

The kinetics observed at 973 K and P = 10 -4 atm were z initially very slow followed by accelerating kinetics. The rate while accelerating was slightly slower than that 94

Table 11 Sulfidation Rate Constants of Fe-4.5Mn-8.8Al-0.36C

v /(g cm “2 sec "I,)

973 1073 1173 ?S2 1

I o — 1 1.87 x 10"6 VO 1 o 1 —

1 4.32 x 10~7 3.69 x 10"6 VO I l 1 CM o o r-~ i o — — X 1 i 7.20 x 10~7 6.88 x 10"7 n i I 1 1 i CM i I o 00 o O LD o i o — — — — X X i 1 i i 4.14 x 10~7 1073 ( ' W3

' DS

3W3

0

W

Figure 18: Linear plots, weight gain vs time, for the sulfidation of the alloy at 95 1073 t

'

W3

'

QS/9W

)

H/WTT

Figure 19: Linear plots, weight gain vs time, for the sulfidation of the alloy at 96 1073 K in Pc = 10-7 atm. 1073 ro LD

fD O I rH o fa 4-> CO ii f0 £

B B B CN LD ^_W3

B

B ; CN O

dw z — ID

i to/ wvj IS 4

CZ Q ** LD o I'D ID O O Q ID o O O ID O O O

TIME (MINI •H 04 rH •H o r •H fa ,3 M-l • i 4-1 4- •H ci M-4 i i 4-> — 4-1 4-1 — — — 4-> tP 5-1 CD 3 cd 5- fd o u a o £ tr> £ C o ^4 (D CO 3 h (d o c O fd 1 ( 1 I I i i

i

r •H o 00 0- — c I CO 04 9 1173 (

*W3

’ OS

* DW )

H

W «d* LD O * x h- X LU — »

i i -P X — 4-1 — ■H -p •H rd o >1 4-1 -P rd (D Xi o 4-t O -P G •H rd 0) ■ CD G O U ■H x g CD p -p r CO i i G > rd -P -H CP

CP CM g •H — co O Q-i C rd U G CD CD U 98 CP G - m

I

r- I i V .

o -p < « G

rH C4 6 — ’ o rd co r-

H w I

.

99

-5 observed at P = 10 atm. The rate constants are shown in b2 Table 11. The linear rate constants were obtained by regression of the appropriate portions of the data on Equation (31).

The sulfidation kinetics observed at 1073 K are shown in figure 18. At 1073 K and P = 10 ^ atm, slow protective b2 behaviour was observed (see figure 19). Parabolic plots of the same data, shown in figure 20, indicates parabolic kinetics. The parabolic rate constant obtained by regression -9 2 -4 of data on Equation (30) was found to be 1.77 x 10 g cm -1 sec

_7 At 1073 K and sulfur partial pressures greater than 10 atm, a short induction period followed by accelerating kinetics were obtained. The duration of the induction period tended to decrease as sulfur partial pressure increased. If the period of rapid reaction during the latter stages of these reactions is approximated by a linear reaction equation, then it is found that the linear rate constants obtained increases as the P increases. This is shown in Table 11. z

Figure 21 shows the sulfidation kinetics obtained at 1173 K. -7 At P = 10 atm, there was initially a slow reaction followed b2 by an acceleration in rate and then fast linear kinetics. At

P = 10 6 atm, the initial slow reaction proceeded at a rate b2 -7 similar to that observed at P = 10 atm, but the final 2 -5 linear rate was faster. At Pq = 10 atm, there was an b2 initial induction period of about 60 minutes followed by fast linear kinetics. At P = 10 atm, there was no induction b2 100

period and the rate immediately proceeded at the most rapid linear rate of all. The induction period decreased as the sulfur partial pressure was increased. -40 »- K) d 1 CN O I (

I 1 & ■P fd £ 4 M) c n CN

4 N1 O 1 — l I

4 - O CM LD o

Figure 22: Plot of -In k-. vs 1/T (K ) for the alloy sulfidised at P = 10 atm. 101 ♦ ------♦ -30 ------‘

CM O I ►

(

'MJNl O I

- O T

CM

Figure 23: Plot of -In k, vs 1/T (K ) for the alloy sulfidised at P = 10 atm 102 103

7.1.1 Temperature Dependence of Sulfidation Rate

Figure 22 and 23 show that, overall, the linear rate constants increase with temperature for constant P -5-4 b2 values of 10 atm and 10 atm. Approximate activation energies for sulfidation were estimated from _ R these Arrhenius plots at P, 10 atm of 100 kJ/mole .-4 and at P, atm of 170 kJ/mole. 104

hO 1 1073

and

K

1173

<•

1 at

sulfidised

alloy

the

for

)

(P

log

vs

k-.

-log

of r\ i Plots

24:

CD Figure I

K) in cd r\. CD i i i i i 1 f>TD01 ■ 105

7.1.2 Pressure Dependence of Sulfidation Rate

Figure 24 shows the effects of sulfur partial pressure on the linear rate constants at 1073K and 1173K. The rate increases as the sulfur partial pressure is increased. The dependence of the rate constants on the sulfur partial pressure was obtained by regression analysis on the double logarithmic representation:

ki 00 ps.

log k7 cc n log P X b2

1 where n is the slope of the log k, versus log P plots. The 1 b2 apparent slope (n), calculated from the results was found to be 1/3 at 1073K and 1/4 at 1173K, thus

1/3 at 1073K ki 00 ps.

1/4 at 1173 ki 00 ps 106

Figure 25: Scale formed on alloy sulfidised at

T = 973 K and P = 10 ^ atm for 7.4 hours b2 (x 1300) . 107

Figure 26: Scale formed on Figure 27: Inner layer of edge of specimen sulfidised at scale formed on edge of T = 973 K and P = 10~7 atm specimen sulfidised at T = z for 4.3 hours (x 130). 973 K and P = 10 7 atm for b2 4.3 hours (x 13 00).

Figure 28: Outermost scale formed on edge of specimen sulfidised at T = 973 K and

Pc = 10“ atm for 4.3 hours 2 (x 1300). 108

Figure 29: Electron probe microanalysis of specimen

sulfidised at T 973 K and P = 10 -7 atm, b2 109

ALLOY . INNER LAYER OUTER LAYER

i/i 40 ■

DISTANCE C/i)

Figure 30: Electron probe microanalysis of specimen

sulfidised at T = 973 K and P = 10 atm . b 2 7.2 Sulfide Scale Morphology and Phase Constitution

7.2.1 T = 973 K, P = 10 8 atm z A uniform layer of fine black scale was formed on the sample sulfidised at 973 K and P = 10 8 atm. The protective film z being extremely thin, was barely resolvable under the microscope (see figure 25). However, in situ X-ray analysis revealed the presence of oc- MnS only (see appendix)

7.2.2 T = 973 K, P = 10_7 atm z - 7 The scale formed at 973 K and P = 10 atm consisted z mainly of a thin compact layer. Severe localised corrosion was observed on the specimen edge, but there appeared to be no tendency for this attack to spread by disrupting the adjacent thin, protective scale (see figures 26 to 28) which covered the flat surfaces of the specimen. X-ray analysis indicated that the scale was composed of oi- MnS, Fe.^ S and small traces of M A^S^ (where M is Fe and/or Mn) .

Electron probe microanalysis of the thin layer of scale -7 formed on specimen faces at 973 K and P = 10 atm is z shown in figure 29. The alloy subsurface region displays a slight depletion of manganese accompanied by a slight increase in iron concentration. No sulfide was detected in the alloy subsurface region. The thin layer of scale was found to consist of oc- MnS.

The distribution of iron, aluminium, manganese and sulfur in the nodule formed at the specimen edge is shown in figure 30 Ill

The alloy subsurface region is strongly depleted of both Mn and Al and the Fe concentration is correspondingly increased.

The light coloured crystalline outer layer was observed to consist mainly of FeS. This layer was enriched in Mn at its outer surface. However, the Mn content decreased to ~ 1 at.

% as the inner layer - outer layer interface was approached from the outer edge. X-ray diffraction analysis showed the presence of oC- MnS, and it seems likely that the outermost Mn-rich region of the scale contains this phase.

There was a slight increase in the Fe and Al concentration near the inner layer - alloy interface with an insignificant concentration of sulfur detected. The inner layer exhibited an uneven distribution of the elements, in particular it was observed that troughs in the manganese profiles tended to coincide with the troughs in the sulfur profile as the aluminium profile rose.

A high iron content within the outer scale layer was apparent as was an aluminium and manganese enrichment in the darker inner layer. However, it was not possible to determine the composition of the inner layer due to the porous nature of the scale. The inner layer is probably a mixture of oc - Mn(Fe)S and M Al-S^ which were detected by in situ

X-ray diffraction analysis. The appearance of this inner layer (figure 28) is consistent with a two phase constitution.

7.2.3 T = 973 K, P = 10~6 atm b2 At 973 K and P = 10 6 atm, a duplex scale was formed, which z 112

Figure 31: Scale formed on Figure 32: Scale formed on alloy sulfidised at T = 973 K alloy sulfidised at T = 973 K and P, , „-6 atm for 4.7 and P = 10 — atm for 4.6 b2 hours (x 130). hours (x 1300).

Figure 33: Tiny nodule

formed on alloy sulfidised

at T = 973 K and P = 10~ b2 atm for 4.6 hours (x 130) . 113

resin inner resin needles ''layer *

DISTANCE (jn)

.Figure 34: Electron microprobe analysis of scale formed

on nodule of specimen sulfidised at 973 K

and P = 10 atm. z COMPOS!TION(WT%) Figure

ALLOY 35 :

Electron on 973

flat K

and

LAYER

surface INNER DISTANCE(

microprobe P b 2 =

of 10

specimen analysis atm. m ) OUTER

sulfidised of

scale LAYER

formed at 114 115

which consisted of a dark inner layer and a lighter outer layer as shown in figures 31 and 32. The outer layer consisted of many discrete, needle-shaped crystals. Tiny nodules were also observed as shown in figure 33. X-ray diffraction analysis showed that this scale consisted of

- MnS, Fe,1 -y S and traces of M Al^S.2 4 .

Results of electron probe microanalysis of the tiny nodules formed on the specimen sulfidised at 973 K and P = 10 ^ atm b2 are shown in figure 34. The outer layer was shown to be of a relatively uniform composition, averaging Fe Al Mn S__ _. This is consistent with a phase constitution 2.7 50.0 of FeS (as confirmed by X-ray diffraction analysis - Appendix A) containing traces of dissolved Mn and Al. From the figure, it can be seen that Fe is enriched in the outer layer of the scale and depleted in the inner layer. The microprobe scans show high concentrations of aluminium and manganese in the inner layer coinciding with troughs in the iron profiles respectively. This suggests the formation of aluminium, manganese and iron sulfides at the inner layer. It is most probably a heterogeneous mixture of oc - MnS and M A^S^ as suggested by the results of in-situ X-ray diffraction analysis of the scale. The porous nature of the inner layer made quantitative assessment of the relative amounts of the sulfides impossible.

The distributions of iron, aluminium, manganese and sulfur in the uniform duplex scale of this specimen are shown in figure 35. The profiles are similar to that shown in figure 34. 116

Figure 36: Scale formed on Figure 37: Scale formed on edge of specimen sulfidised at edge of specimen sulfidised at T = 973 K and P = 10 ^ atm T = 973 K and P = 10 atm b2 b2 for 4.8 hours (x 130). for 4.8 hours (x 520).

Figure 38: Scale formed on Figure 39: Scale formed on flat surface of specimen flat surface of specimen sulfidised at T = 973 K and sulfidised at T = 973 K and -5 P = 10 atm for 4.8 hours PQ = 10 atm for 4.8 hours b2 b2 (x 130) . (x 130) 117

^

at formed

OUTER scale LAYER

of

sulfidised

. RESIN atm analysis specimen

10 of

= 2 b

P microprobe

surface INNER

LAYER and

K DISTANCE^) flat

973 on Electron

: 40 ALLOY

Figure (%’ 1 /A } N ofll S OdH 0 3 118

The outer layer was again shown to be composed of mainly FeS with traces of Mn and Al (derived from the composition, Fe,, _ Al, „ IVhi S . _ . However, the outermost part of the outer layer contained a slightly increased amount of

Mn. The dark inner layer exhibited an uneven distribution of elements similar to that of figure 34.

From figure 34 and 35, it was evident that the alloy sub­ surface area is strongly depleted of both Al and Mn whereas the inner layer of scale is strongly enriched in both Al and Mn.

7.2.4 T = 973 K, P = 10~5 atm z As shown in figure 36 and 37, the thick scale formed on the specimen sulfidised at 973 K and P - 10 -5 atm was duplex in z nature, with a thicker darker inner layer and a lighter coloured outer layer with fine whiskers at the outermost edge. When observed very carefully, a very fine layer of scale between the darker inner layer and the lighter coloured outer layer was noted. Nodules as shown in figure 38 were observed on the flat surfaces of the specimen. X-ray diffraction analysis indicated that Fe_ S, oc - MnS and M Al~S, are l-y z 4 present. Shallow internal sulfidation was also observed as shown in figure 39.

Figure 40 displays the distribution of the metallic elements and sulfur in the scale formed on the specimen sulfidised at _5 973 K and P = 10 atm, as obtained from microprobe analysis, z The microprobe scans reveal a high concentration of Fe in the outer layer. The X-ray intensities for all the elements show 119

Figure 41; Scale formed on Figure 42: Scale formed on edge of specimen sulfidised at specimen sulfidised at

T = 973 K and P = 10~4 atm T = 973 K and P = lO-4 atm b2 b2 for 4.9 hours (x 130). for 4.9 hours (x 520).

Figure 43: Scale formed on Figure 44: Scale formed on alloy sulfidised at T = 973 K alloy sulfidised at T = 973 K -4 -4 and P =10 atm for 4.9 and P = 10 atm for 4.9 hours 2 b2 hours (x 520). (x 320). 120

Figure 45: Scale formed on specimen sulfidised at

T = 1073 K and P = 10 ^ atm for 5 4 hours z (x 1300) . 121

a thick scale layer with a relatively uniform composition of

Fe^ _ Alrt „ Mn „ SAn „ which is consistent with a layer made up mainly of Fe (Mn)S with traces of dissolved Al.

Within the inner layer, the "peak" of Al and S appeared to coincide with the "trough" of Fe content. The inner layer

was found to have a composition Fe2Q A-*-24 7 Mn2 2 S53 1 which is approximately consistent with a mixture of

FeA^S^ and Fe(Mn)S.

Again a significant depletion of aluminium and manganese was evident at the alloy subsurface.

7.2.5 T = 973 K, P = 10~4 atm b2 As shown in figures 41 to 44, a fine protective scale was observed with nodules formed mainly at the specimen edges. The nodules formed at the corners of the specimen consisted of a porous multiphase inner layer and a lighter coloured

outer layer. The presence of Fe^ S, oc - MnS and traces of MAl2S^ were detected by in situ X-ray diffraction analysis.

Small needle-like protrusions into the adjacent alloy were -7 observed for the samples sulfidised at 973 K and ^ 10

7.2.6 T = 1073 K, P = 10 7 atm b2 A fine adherent single layer of grey scale was observed on

the specimen sulfidised at 1073 K and P = 10 ^ atm. X-ray z diffraction analysis revealed the presence of A^S^ only. Figure 45 illustrates the metallographic cross section of 122

Scale formed on Figure 47: Scale formed on

edge of specimen sulfidised at edge of specimen sulfidised at -5 T = 1073 K and P = 10*"5 atm T = 1073 K and P = 10 atm z z for 5.9 hours (x 20). for 5.9 hours (x 500).

Figure 48: Internal layer of

nodule formed on edge of specimen sulfidised at

T = 1073 K and P = 10 ^ atm b2 for 5.9 hours (x 500). 123

Figure 49: Tiny nodule formed on flat surface of specimen

sulfidised at T = 1078 K and P = 10 6 atm b2 for 7.8 hours (x 230).

Figure 50: Tiny nodules formed on flat surface of specimen

sulfidised at T = 1078 K and P = 10 ^ atm U for 7.8 hours (x 200). 124

Figure 51: Nodule formed on Figure 52: Magnified view of flat surface of specimen nodule adjacent to protective sulfidised at T = 1073 K and layer of scale on flat surface PQ = 10 4 atm for 1.9 hours of specimen sulfidised at b2 (x 130). T = 1073 K and P = 10 4 atm b2 for 1.9 hours (x 520).

Figure 53: Specimen sulfidised

at T = 1073 K and P = 10~4 b2 atm for 1.9 hours (x 520). 125

this scale. Since the expected oc - MnS was not detected in the scale, it appeared that the oc- MnS layer might have fallen off. X-ray diffraction analysis of a replicate sample revealed the presence of and oc - MnS.

7.2.7 T = 1073 K, P = 10~6 atm and P = 10_5 atm S2 z -5 In the specimens sulfidised at 1073 K and at P = 10 atm, b2 a continuous layer of nodular growth all around the specimen edges was observed as shown in figure 46. Magnified micrographs of the nodule formed at the specimen edge is shown in figure 47 and 48. A few tiny nodules on the flat surfaces were also observed at these pressures as shown in figure 49 and 50. Needle-like protrusions into the alloy are observed under the tiny nodules. A fine protective layer covers the remaining surface of the specimen.

7.2.8 T = 1073 K, P = 10~4 atm b2 At P = 10 -4 atm, larger nodules were formed on the flat z surfaces as well as on the specimen edges. The nodules had a complex and fragile structure as shown in figure 51 to 53. The outer scale was silver grey in colour, with a highly crystalline appearance. The thick multiphase inner layer of the scale exhibits a fine fibrous structure with fibres orientated normal to the alloy surface. A fine protective layer of scale covers the remaining surface. X-ray diffraction analysis detected oc - MnS, Fe„ S and MA1„S„. l-y 2 4

Electron probe microanalysis of the nodules formed on the 126

ALLOY INNER LAYER OUTER LAYER

40 8 0 120 160 200 DISTANCE^)

Figure 54; Electron microprobe analysis of a nodule formed on specimen sulfidised at 1073 K and

PQ = 10 atm. 2 127

Figure 55: Scale formeid on edge of specimen sulfidised -7 at T = 1173! K and P - 10 atm for b2 4.0 hours ((x 130)

\\

9 m 1 Figure 56: Needle-like Figure 57: Internal layer of protrusions into alloy scale formed on edge of sulfidised at T = 1173 K and specimen sulfidised at _7 P = 10 atm for 4.0 hours T = 1173 K and P = 10 ^ atm b2 b2 (x 520) . for 4.0 hours (x 520). 128

Thin layer

Figure 58: Scale formed on flat surface of specimen sulfidised at T = 1173 K and atm for 4.0 hours

(x 130).

Figure 59: Scale formed on specimen sulfidised at

T = 1173 K and P = 10 ^ atm b2 for 4.0 hours (x 130) 129

Figure 60; Internal layer of scale formed on specimen sulfidised at T = 1173 K and P = 10 ^ atm for 3.9 hours b2 (x 520) .

Figure 61: External layer of scale formed on specimen sulfidised at T = 1173 K and

P = 10 ^ atm for 3.9 hours z (x 520) . 130

Figure 62: Scale formed on alloy sulfidised at T - 1173 K and P = 10 ^ atm for 2.7 b2 hours (x 65).

Figure 63: Scale formed on

specimen sulfidised at

T = 1173 K and P, 10 atm

for 2.7 hours (x 320)

Figure 64: Scale formed on specimen sulfidised at

T = 1173 K and P = 10 ^ atm b2 for 2.7 hours (x 130). 131

Figure 65: Scale formed on Figure 66: External layer of

edge of specimen sulfidised at scale formed on specimen

T = 1173 K and P = 10~4 atm sulfidised at T = 1173 K and b2 -4 for 1 hour (x 65). P = 10 atm for 1 hour b2 (x 130) .

Figure 67: Internal layer of

scale formed on specimen

sulfidised at T = 1173 K and

PQ - 10 4 atm for 1 hour b2 (x 320) . 132

-4 specimen sulfidised at T = 1073 K and P = 10 atm is shown b2 in figure 54. The outer region of the scale constituted by the "whiskerlike" growth was observed to consist mainly of

Fe(Mn)S although a trace of Al (0.2 at. %) was evident in this region. Again the heterogeneous character of the inner

scale was evident. The inner layer was enriched in aluminium and contained quite a high concentration of Fe. The aluminium "trough" region of the inner layer revealed a composition Fe37 Al Mn S54 5 which is approximately consistent with a mixture of FeA^S^ and Fe(Mn)S. There was noticeable depletion in the Al and Mn content in the under­ lying alloy substrate.

7.2.9 T = 1173 K

At 1173 K, the nodules had coalesced to form a relatively uniform, stratified scale with a rather flaky outer layer. Figures 55 to 67 illustrates the topographical features of the scale present on the specimen sulfidised at 1173 K. At 1173 K, the scale generally consisted of a thin crystalline external layer and a thicker multiphase inner layer of porous looking material. A very fine film of scale was also observed between the internal and external layer of scale

(see figures 58, 59, 63 and 64). Small needle-like protrusions into the adjacent alloy were observed for the samples sulfidised at 1173 K. However, the needle-like protrusions into the alloy were observed to be deeper at this temperature than those observed at the lower temperatures.

Partial spallation of the outer scale was always observed 133

Ul C£ LU uo zn 5: at

sulfidised

specimen

of

edge

of

analysis

microprobe

Electron

: 68

Figure

°°(%“iM)N0ilIS0dW0D 134

during cooling from reaction temperature (1173 K) as shown in figure 64 and 67. Even when thermal shock was decreased by slow furnace cooling, partial spallation of the scales was observed. Separation occurred at the external scale - internal porous scale interface but not at the sulfide - metal interface.

oc - MnS, Fe„ S and traces of MAlnS, were detected by X-ray 1-y z 4 diffraction analyais in all the specimens sulfidised at 1173 K except the one sulfidised at P -6 atm where only Fe„ S 1 — v and oc - MnS were detected.

Electron probe microanalysis of the scales formed on the -4 specimen sulfidised at T = 1173 K and P = 10 atm is shown z in figure 68. In the outer region of the scale, made up of the "whiskerlike" growth, the observed sulfur concentrations were quite low with the iron concentration being correspondingly high. This part of the scale was mainly composed of FeS.

As shown in figure 68, the porous looking heterogeneous layer which underlies the "whiskerlike" growth is composed mainly of Fe(Mn)S and some FeA^S^ (derived from the composition,

Fe47 1 8 Mni 4 ^49 7 °^)Serve(^ 136 yum away from the alloy surface) .

The darker innermost region of this scale showed an enrichment of Al and Mn. The inner layer penetrated into the alloy as needle-like protrusions to a depth of approximately 15 yum. The very uniform distribution of the internal needles appeared to be responsible for improved scale adherance. The region 135

of the internal needles was considerably depleted in manganese and aluminium and enriched in iron. 136

8.00 DISCUSSION

8.1 Introduction

Because of the complexity of the rate curves, it seems unlikely that a single mechanism is operative throughout the reaction. Attention will therefore be focussed on the development of sulfide structure and morphology to detect any changes which may indicate changes in growth mechanism.

The observed reaction kinetics were of two types. On the one hand, when the combination of temperature and sulfur partial pressure was not too severe, protective scaling kinetics resulted. Thus sulfidation kinetics were _ 7 approximately parabolic at T = 973 K when P ^10 atm, b2 -7 and at T = 1073 K when P = 10 atm. On the other hand, b2 initially accelerating reaction rates followed by fast linear kinetics were observed at high temperatures and sulfur partial pressures. Behaviour of a somewhat intermediate type was observed at T = 973 K and P = 10 ^ atm. b2

Protective kinetics accompanied the formation of thin, uniform scales rich in manganese and, in one instance, aluminium also. Fast linear kinetics were observed when large sulfide nodules grew with the formation of a thick, two phase reaction product layer at their base. Discussion will be directed in turn to these two aspects of the sulfidation reaction. 137

8.2 Protective Scaling Kinetics Of all the specimens sulfidised, the one reacted at the lowest temperature and sulfur pressures tested (T = 973 K, P =10 atm, 10 atm), possessed the highest sulfidation S2 resistance. This appears to be due to the continuous adherent product layer of fine grained black sulfide formed which consisted of oc - MnS.

_ 7 At 1073 K and P =10 atm, a protective scale which was b2 found to consist of A^S^ under an oc - MnS layer grew according to parabolic kinetics.

The presence of oc - MnS adjacent to the alloy surface is in accordance with the fact that oC - MnS is the most stable of the three possible sulfides as shown in Table 8. oc - MnS is the most stable sulfide and therefore the preferred reaction product and therefore can be expected adjacent to the alloy surface. Apparently at 973 K, the material is protective, i.e. does not permit the diffusion of Fe to form FeS on top. Previous research on the sulfidation of manganese (5) and iron is relevant to this observation and will be compared with the findings of this investigation.

Elrefaie and Smeltzer v ’ investigated the sulfidation _ g behaviour of pure Mn in atmospheres (10 atm p ^10"4 atm) at temperatures between 1073 and 1273 K and S2 observed that the reaction kinetics obeyed the parabolic rate law. When manganese was sulfidised at T = 1073 K and -7 -9 2 P = 10 atm, a parabolic rate constant of 1.2 x 10 g cm -4 s 1 was obtained. In the present investigation, 138

Table 12: Parabolic Rate Constants of Iron^1^ and

Fe-4.5Mn-8.8Al-0.36C Alloy

, 2 -4 -1. kp (g cm s )

Fe Fe-4.5Mn-8.8A1-0.36C

T = 1073 K 7 Q ~j 2.3 x 10 1.8 x 10 P = 10 atm S2

T = 973 K „ -7 „ -7 -10 P_ = 10 atm 1.7 x 10 7.8 x 10 S2 P =10 atm 1.0 x 10~7 4.9 x 10"10 b2 139

when the alloy was sulfidised under the same temperature and sulfur partial pressure, a parabolic rate constant of

1.8 x 10 -9 g 2 cm -4 s -1 was obtained. There appears to be fairly good agreement but, of course, A^S^ was also present in the scale investigated.

When Elrefaie and Smeltzer's data are extrapolated to -10 2 -4 T = 973 K, a parabolic rate constant, kp = 7 x 10 g cm s -1 is obtained for P = 10 -7 atm and a value of kp = b2 4 x 10 —10 g 2 cm —4 s —1 for P = 10 —8 atm. In comparison, b2 alloy sulfidation at 973 K gave a parabolic rate constant of 7.8 x 10 ^ g^ cm 4 s 4 for the specimen sulfidised at -7 -10 P = 10 atm and a parabolic rate constant of 4.9 x 10 b2 _8 g2 cm 4 s 1 for the specimen sulfidised at P = 10 atm. 2 There is seen to be excellent agreement between the two sets of results at 973 K. It is concluded therefore that alloy sulfidation under these conditions is controlled by the rate of growth of an external MnS scale.

Previous research on the sulfidation properties of iron will now be compared with the presently reported results on

Fe-4.5Mn-8.8Al-0.36C (see Table 12). It can be seen from Table 12 that the sulfidation rate of Fe-4.5Mn-8.8Al-0.36C alloy is approximately three orders of magnitude less than that of pure iron. This comparison is of course limited to conditions under which parabolic kinetics are observed, i.e.

T = 1073 K when P = 10 ^ atm and T = 973 K when P ^ 10 ^ b2 2 atm.

The parabolic rates of Fe-4.5Mn-8.8A1-0.36C alloy are 140

approximately the same order of magnitude as the rates of pure manganese, and approximately three orders of magnitude less than the parabolic sulfidation rate of pure iron. The _ 7 parabolic behaviour observed at 973 K and P

From a thermodynamic calculation, it can be shown that for

MnS to be destabilized with respect to Al2S2 through the reaction

3 MnS + 2A1 = Al^ + 3 Mn 3 2 -3 would require the activity ratio (a ) / (a -j_) ^ 5.7 x 10 at T = 1073 K. For example if a^n = 0.010, then Al2S2 would 141

Figure 69: Activity Gradients of Mn and Al

Alloy 142

Table 13: Parabolic Rate Constants for Sulfidation at

T = 973 K

4 -1, kp (g^ cm s )

P = 10 8 atm P =10^ atm b2 b2

Fe-4.5Mn-8.8A1-0.36C 4.9 x 10-10 7.8 x 10"10

Fe-25Mn-5Al 1.0 x 10~9

Fe-26.6Cr 7.9 x 10"8 8.9 x 10~8

Fe-28Mn 1.1 x 10"8 1.7 x 10'8

Table 14: Parabolic Rate Constants at T = 1073 K and

P = 10 ^ atm b2

2 -4 -1 kp (g cm s )

-9 Fe-4.5Mn-8.8A1-0.36C 1.8 x 10

Fe-25Mn-5Al 3.0 x 10"10

-8 Fe-2 8Mn 3.0 x 10

^ ^ „ -7 Fe-26.6Cr 3.2 x 10 143

form if a 0.013. The concentration gradient of Mn can be illustrated by figure 69. From the above discussion, it can be expected that alloy depletion in Mn could lead to

Al2S3 formation at the base of the scale. Therefore the formation of an product layer under an ex - MnS layer is explicable within a framework of local equilibrium. Evidently this scale functions as an effective barrier to Fe diffusion.

Previous investigations on the sulfidation of Fe-25Mn-5Al^^ and Fe-28Mnv ; will be compared with the kinetic results of this investigation. Table 13 demonstrates that Fe-4.5Mn-8.8 A1-0.36C alloy exhibits superior sulfur corrosion resistance at 973 K and P = 10 ® atm and P = 10 ^ atm. At 1073 K _7 2 2 and P = 10 atm, the alloy was also found to be superior z to Fe-28Mn by approximately one order of magnitude, as shown in Table 14. However, under these conditions, this alloy was found to be less sulfur resistant that Fe-28Mn-5Al by approximately one order of magnitude.

Fe-4.5Mn-8.8Al-0.36C alloy exhibits more corrosion resistance than Fe-26.6Cr^^ at T = 973 K and P = 10 ^ atm and 10 -7 atm, this effect being of the order of two2 -7 magnitudes. At T = 1073 K and P = 10 atm, the corrosion z resistance of the presently studied alloy was found to be superior by one order of magnitude.

The sulfur corrosion resistance of the presently studied alloy was found to be approximately two orders of magnitude better than the resistance for the stainless steels with the 144

composition, Fe-Cr-8Ni with 4, 12 and 22 wt % Cr^7^ but

this effect is confined to the lower sulfur partial pressures and temperatures, where the Fe-Al-Mn alloy reacts with parabolic kinetics. The parabolic rate constants for -7 Fe-Cr-8Ni with 4, 12, 22 wt % Cr varies between 10 and

10 ** g2 cm 4 s l at P = 3.2 x 10 ^ atm and T = 1023 K. In b2 comparison, the parabolic rate constant for the presently -10 2 -4 -1 studied alloy is only 4.87 x 10 g cm s at T = 973 K and P = 10 — 8 atm. The parabolic rate constant for the b2 -7 specimen sulfidised at T = 1073 K and P = 10 atm is b2 1.77x10i -- 1n -9 g 2-4-1cm s

It may be concluded that the scale formed on Fe-4.5Mn-8.8Al- 0.36C at these relatively low temperatures and sulfur partial pressures, is more protective than the scale formed on austenitic stainless steel. Under these circumstances, Mn is thus a more beneficial alloy additive than is Cr. This conclusion has already been reached for high Mn content

/ r \ alloys1 . The significance of the present result is that quite a low level of Mn has been shown to be sufficient for protection against sulfur corrosion. It is also concluded that under these reaction conditions, the presently studied

alloy is more resistant to sulfidation than Fe-Mn binary

alloys. This must reflect the effect of the alloy Al content. Whilst the Al plays a part in scale formation at 1073 K and P = 10 7 atm, its role in the alloy sulfidation b2 behaviour at 973 K is unclear. At both temperatures, protection is achieved by preventing the formation of

Fe-rich sulfides. Because low levels of Mn are thought (4) 145

not to have such an effect, it is presumably to be attributed to the presence of Al. The way in which Al performs this function is not revealed by the present results. 146

A12S 3

INDUCTION PERIOD

ALLOY

ACCELERATING KINETICS

Mn, Al Sulfides

ALLOY

LINEAR KINETICS

FIGURE 70 PROPOSED MECHANISM OF NODULE GROWTH 147

8.3 Fast Nodule Growth Kinetics

- 5 -4 At 973 K and P values of 10 atm and 10 atm, 1073 K and S 2 10 6 ^ Pc <1 10 ^ atm and at 1173 K and 10 P 10 ^ atm, S2 2 breakdown of the protective oC -MnS scale was observed with the formation of large sulfide nodules which possessed a thick, two-phase reaction product layer at their base. This was associated with initially accelerating reaction rates followed by fast linear kinetics.

The proposed mechanism of nodule growth will now be illustrated

(see figure 70). During the induction period, a fine layer of scale is formed containing oc -MnS and A^S^* The activity gradient of Mn and Al are shown schematically in figure 69.

After an initial "induction period", there is a sudden increase in rate known as "break-through". "Break-through" seems to occur at decreasing times as the sulfur pressure is increased. The very early stages of "break-through" on this alloy are difficult to detect metallographically because, once it occurs, it proceeds extremely rapidly. "Break-through" is attributed to mechanical failure of the protective scale which allows gas access to the underlying metal leading to nucleation and rapid growth of the nodules. Figure 53 illustrates a typical scale nodule adjacent to a region of protective scale, part of which was adherent and the rest non-adherent.

Because the underlying alloy is depleted of Mn and Al repassivation by the formation of MnS or A^S^ is impossible.

Instead Fe can react, leading to the formation of the 148

multiphase scale containing sulfides of Fe, Mn and Al (see

figure 70 (b)). This is the nucleation event. The nodule

material is porous and has an effective large molar volume.

Its continued growth leads to progressive disruption of the

formerly protective scale. As the nodules expand laterally

and thicken , the overall rate of sulfur uptake accelerates as

seen in the kinetic behaviour at 1073 K.

The causes of the initial mechanical failure of the protective

scale have not been precisely established as pointed out

previously, (25,27,30,31,32) sca2e failure is determined by

the adhesion of scale and alloy, the development of stress in

each and the ability of each to deform and relieve the stresses.

Failure is a complex function of alloy composition, surface

finish, sulfidation temperature, scale thickness, atmosphere

and specimen geometry. Scale failure is easier the sharper

the specimen angles and the higher the temperature and sulfur pressure.

The scales which become detached from the alloy assume a wrinkled appearance indicating they have been under compressive stresses and possibly this is the chief cause of

failure. The remains of the initial fine layer of protective

scale can be seen in figure 58, 63 and 64. Figure 59

illustrates the breakthrough in the initial fine protective

layer which seems to have occurred under compressive stresses.

The superior adhesion of scales developed at 973 K may be due

to less stress development related to the more stoichiometric

structure and lower formation rate. On the other hand, they TABLE 15: Linear Rate Constants obtained by sulfidation of Fe-4.5Mn-8.8Al and Fe-25Mn-5Al t CM H 1 i M tT> O & o 0 CD rH P-t CA CM rH ro ro cn o 1 1 7 3K rH rH rH CM i 00 CM rH m m oo m m "sT rH m 00 oo m m 00 oo CM a a C C a m mil P a < C — C

• • • • • • 1

vo t Ifl

till I

x

o CT> — o < <

o \ VO i h — — — x x 1

• 1 1 i

CM I vo vo LD

rHo

r- ro rH o o rH i

— x x

• i

vo t"* m mm mm U0 x o i i I I I i — o VO r o ro cn «H o CM cm rH

r r" o — CM CTi CO r- i o

— — —

XX xx XX xx x x 1

1 • •

1 1 i

r- h h I i i i x o

— o CM t I 1 1 rH o ^ Ox o 1 rH cm iH o oo m — — — 1 1 1 1 149 150

may still originally be heavily stressed but are merely more adherent or more plastic. Probably the adhesion is partly improved under these conditions because fewer vacancies are arriving at the alloy/sulfide interface allowing vacancy annihilation at this interface rather than void formation.

In essence, the model consists of the proposition that islands of sulfide nucleate and their growth is limited by the rate of incorporation of sulfur at their periphery, possibly through cracks and pores, the sulfur deriving from the atmosphere during early stages of exposure. Thus as the nodules expand, the overall rate of sulfur uptake accelerates.

As the nodules impinge on one another, the nodules continue to thicken but cease to expand laterally, and so the overall sulfidation rate becomes constant as shown in the results obtained at 1173 K. This behaviour is illustrated in figure 70 (c) .

At 1173 K, partial spallation of the outer scale was always observed. Separation occurred within the external scale but not at the sulfide-metal interface as shown in figure 62. The very uniform distribution of needle-like protrusions into the alloy was probably responsible for the improved adherance of the inner scale.

Previous research on the sulfidation of Fe-25Mn-5Al^^ and

Fe-25Mn-20Al^ will be compared with the kinetic results obtained in the present investigation. Table 15 shows the superior sulfur corrosion resistance of Fe-4.5Mn-8.8Al-0.36C alloy at all the sulfidation conditions compared except at 151

1073 K and P = 10 ^ atm where Fe-25Mn-5Al alloy was found to b2 sulfidise approximately one order of magnitude slower. At 973 K and P = 10 b, 10 ^ atm and when T = 1073 K and b2 P = 10 -5 atm, the presently studied alloy was found to react b2 approximately two orders of magnitude slower than Fe-25Mn-5Al.

When sulfidised at 1173 K and 10 ^ XT PQ ^ 10 ^ atm, the alloy b2 ^ under investigation was found to be one order of magnitude superior in sulfidation resistance.

Fe-25Mn-20Al alloy suifidised according to sigmoidal kinetics at 1073 K and P = 10 atm and 10 ^ atm. When the b2 period of rapid sulfidation was approximated by a linear rate equation, , the linear rate constant was found to be 3 x 10 6 g cm 2s and 1 x 10 ^ g cm 2s ^ at P = 10 4 atm -5 b2 -4 and 10 atm respectively. At 1073 K and P of 10 atm, the b2 sulfidation linear rate constant of the presently studied alloy -4 (k^ = 2.1 x 10 ) was found to be two orders of magnitude worse than the sulfidation rate of Fe-25Mn-20Al alloy. However at P of 10 -5 atm, the sulfidation resistance (kn = 6.9 x 10 -7 b2 1 g cm -2 s -1 ) of the alloy studied was two orders of magnitude superior to Fe-25Mn-20Al alloy. 152

9.00 CONCLUSION

Sulfidation kinetics and development of sulfide scales on an

Fe-4.5Mn-8.8Al-0.36C alloy were studied at 973, 1073 and 1173 K in H2S-H2 atmospheres at sulfur pressures 10 -4 atm^\ P^ ^

10 — 8 atm. Two types of reaction kinetics were observed.^

Reaction kinetics were protective and slow when the combination of temperature and sulfur partial pressure was not too severe. At higher temperatures and sulfur activities, initially accelerating reaction followed by fast linear kinetics were observed. Behaviour of a somewhat intermediate __ g type was observed at T = 973 K and P = 10 atm. b2 The alloy was found to sulfidise according to parabolic kinetics at T = 973 K when P = 10 ^ atm and 10 ^ atm and at T = 1073 K when P = 10 -7 atm.2 The protective scale formed b2 was uniform, thin and rich in manganese, and in one instance, aluminium also. From the parabolic kinetic results and growth of compact scales, the reaction appeared to conform to Wagner's Theory of Oxidation Rate, and thus under these conditions, the growth of the ot - MnS scale could have been controlled by solid state diffusion of Mn 2+ ions. It can be concluded that

Fe does not participate in the reaction.

Formation of an Al2S3 scale underlying oc - MnS scale was thought to be due to the depletion in Mn.

From a comparison of the parabolic rate constants, it was concluded that Mn is a more beneficial alloy additive than Cr 153

under the reaction conditions specified. This finding was in agreement with previous investigations by Smeltzer et al^ .

It was also found that under these conditions, the presently studied alloy was more resistant to sulfidation than Fe-Mn binary alloys and because low levels of Mn have been thought (4) not to have such an effect, the excellent sulfidation resistance was thought to be due in part to the presence of Al.

At T = 973 K when P = 10 ^ atm and 10 4 atm and at T = 1073 K b2 when 10 P

/ -4 ^ z \ 10 atm, initially accelerating reaction rates followed by fast linear kinetics were observed. This pattern of behaviour was due to mechanical failure of the initially formed protective scales allowing gas access to the underlying metal which leads to the nucleation and rapid growth of nodules. The nodules were found to be composed of Fe(Mn)S and a two phase reaction product layer at their base which was found to be FeCMnjAl^S^ plus Mn(Fe)S. This material was voluminous and appeared to be of high porosity. 154

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(1983) . 80. Flahaut, J., Ann. de Chim., 7 : 632 (1952). PART C - APPENDIX 160

Specimen sulfidised at T = 973 K, P 10 ® atm S2

Sample DC ~ Unsulfidised Sample I/I (%) I/I (%) I/I (%)

3.043 3 . 015

2.600 2.612

2.254 2.266

2.042 2.053

1.838 1.847

1.545 1.575

1.180 1.185

1.168 1.168 161

-7 Specimen sulfidised at T = 973 K, P - 10 atm. b2

(80) (80) Sample FeS

1.023 10.4

1.039 0.1 1.048 10 1.034

1.065 0.3 1.065 7 1.0662 16 1.051 1.0614

1.124 0.3 1.125 7

1.166 0.3 1.167 7 1.168 20 1.1641

1.182 13.9 1.174 10 1.1703

1.210 0.1 1.291 10 1.263 1.2059

1.308 0.2 1.321 20 1.303

1.402 0.1 1.433 20

1.449 0.3 1.442 10 1.487

1.507 0.5 1.490 10 1.509 20 1.5063

1.571 0.2 1.636 10 1.575 6

1.692 0.1 1.720 40 1.653 1.672

1.782 0.2 1.769 7 1.754 1.769

1.847 9.4 1.847 50 1.818 1.826

2.049 100.0 2.064 100 2.071

2.113 0.2 2.160

2.272 0.3 2.255 10 2.289

2.424 0.1 2.415

2.511 0.1 2.576

2.617 0.1 2.640 50 2.612 100 2.683 2.607

2.721 0.2 2.87 10 2.879 2.715

2.911 0.6 2.98 10 2.977 2.916 162

Sample FeS o

I/I. I/It 1/11 d(A) (%) d(A) (%) d(A) (%) d (A) d (A)

3.029 0.2 3.015 14 3.018

3.170 0.2 3.107 3171

3.269 0.1 3.44 7

4.061 0.3 3.980 4.012

5.722 0.1 5.74 20 163

Specimen sulfidised at T =; 973 K, PQ = 10 6 atm b2

Sample Fes oc - MnS FeAl2S4 MnAl2S4 I/I. i/i, I/I» d(A) (%) d(A) (%) d(A) (%) d (A) d (A)

1.017 19 1.034

1.055 9 1.048 10 1.0662 16 1.051 1.0614

1.109 6 1.105 15

1.178 10 1.174 10 1.1682 20 1.170

1.231 1 1.222

1.289 4 1.291 10 1.263

1.333 19 1.321 20 1.340

1.442 9 1.442 10 1.460

1.496 6 1.490 10 1.509 20 1.487 1.506

1.731 100 1.720 40 1.754 1.724

1.828 12 1.847 50 1.818 1.841

1.933 6 1.909 7

2.045 61 2.064 100 2.096 • 2.080 19 2.071

2.578 11 2.576

2.652 48 2.640 50 2.612 100 2.683 2.607

2.988 56 2.980 40 3.015 14 2.977 2.916 L i 164

Specimen sulfidised at T = 973 K, Pc = 10 ^ atm b2

Sample FeS i MnS FeAl2S4 MnAl2S4

I/I, I/I, 0 I/I,

C d (A) (%) d (A) (%) (%) d (A) d (A)

1.020 14 1.034

1.052 7 1.048 10 1.0662 16 1.051 1.0614

1.110 5 1.105 15

1.149 13 1.125 7 1.1682 20 1.1641

1.181 55 1.174 10 1.1703

1.247 3 1.263 1.2215

1.282 4 1.291 10 1.274

1.331 10 1.321 20

1.449 12 1.442 10 1.460

1.486 6 1.490 10 1.487

1.526 8 1.509 20 1.523

1.563 4 1.606 7 1.575 6 1.569

1.619 3 1.636 10 1.634

1.734 100 1.720 40 1.724

1.7508 6 1.769 7 1.754 1.769

1.826 49 1.847 50 1.818 1.841

1.939 7

2.049 62 2.006

2.080 89 2.064 100 2.071 2.096

2.238 8 2.255 10

2.327 9 2.373

2.585 59 2.576

2.667 74 2.640 50 2.612 100 2.683

2.998 27 2.980 40 3.015 14 2.977 3.018 165

-4 Specimen sulfidised at T = 973 K, PQ = 10 atm. ^2

Sample FeS OC ~ MnS FeAl7S4 MnAl^S^ I/I. I/I. I/I. O o o o p d(A) (%) a (A) (%) a (a) (%) d (A) d (A)

1.054 10 1.048 10 1.0662 16 1.051 1.064

1.108 12 1.101 10

1.182 11 1.174 10 1.168 20 1.195

1.295 4 1.291 10 1.263 1.303

1.331 18 1.321 20 1.340

1.376 5 1.384

1.443 16 1.442 10 1.436

1.492 4 1.490 10 1.509 20 1.487

1.603 6 1.606 7 1.575 6 1.653

1.732 76 1.720 40 1.754

1.824 19 1.847 50 1.818 1.826

1.935 8 1.909 7

2.049 80 2.006

2.085 70 2.064 100 2.071 2.096

2.578 28 2.612 100 2.576 2.607

2.667 77 2.64 50 2.683 2.715

3.011 100 2.98 40 2.977 3.018 166

-7 Specimen sulfidised at T = 1073 K, Pc = 10 atm. b2

Sample A12 s3 o o d(A) i/i, <%) d (A) i/i, <%)

1.337 8

1.399 44 1.37 20

1.536 14 1.50 27

1.568 14 1.56 33

1.665 16 1.60 33

1.693 16 1.69 33

1.875 18 1.86 66

2.017 10 1.99 20

2.099 16 2.07 33

2.243 8 2.37 13

2.487 10 2.47 13

2.763 2 2.82 83

3.129 100 2.95 27

3.256 32 3.22 100

3.580 14 3.47 13

4.720 16 4.70 23 167

Specimen sulfidised at T 1073 K, P atm b2 (Replicate sample)

Sample oc -

5.210

4.015

3.712

3.470 3.47

3.235 3.22

3.000 3.015 2.95

2.705 2.82

2.599 26.1 2.612 2.54

2.252

2.036 2.07

1.963 1.99

1.839 11.3 1.847 1.86

1.503 1.509 1.50

1.301 1.306 168

Specimen sulfidised at T = 1073 K, P atm.

Sample FeS (XL - MnS MnAl2S4 FeAl2S4 I/I, I/I, I/I, d (A) (%) d (A) (%) d(A) (%) d (A) d (A)

2.981 36 2.980 40 2.997

2.65 18 2.640 50 2.683

2.612 20 2.612 100

2.040 100 2.064 100 2.096 2.071

1.917 7 1.909 7 2.006

1.843 17 1.847 50 1.841

1.722 37 1.720 40

1.509 2 1.509 20 1.506

1.433 17 1.433 20 1.4336

1.329 6 1.321 20 1.319

1.222 9 1.291 10 1.222

1.180 14 1.174 10 1.1682 20 1.170

1.065 3 1.065 7 1.0662 16 1.061 1.051

1.022 13 1.048 10 1.034 169

Specimen sulfidised at T = 1073 K, P 10 atm. b2

Sample FeS oc - MnS MnAl2S4 FeAl2S4

I/I, I/I, 1 /!» d(A) (%) d(A) (%) d (A) (%) d (A) d (A)

4.054 3 4.012

3.177 3 3.171 3.107

3.010 1 2.980 40 3.015 14 3.018 2.977

2.743 1 2.870 10 2.715

2.604 12 2.640 50 2.612 100 2.607 2.576

2.267 1 2.255 10 2.280

2.046 100 2.064 100 2.096 2.071

1.840 6 1.909 7 1.847 50 1.841 1.818

1.721 3 1.720 40 1.724

1.504 2 1.509 20 1.506

1.367 < 1 1.321 20 1.383

1.301 < 1 1.291 10 1.306 8 1.303 1,181 9 1.167 7 1.168 20 1.195 170

Specimen sulfidised at T = 1073 K, P atm. b2

Sample FeS oC - MnS FeAl2S4 MnAl2S4 I/I. I/I, I/I, d(A) (%) d(A) (%) d(A) (%) d (A) d (A)

3.160 2 3.107 3.171

3.000 4 2.980 40 3.015 14 2.977 3.018

2.642 4 2.640 50 2.6 83

2.604 18 2.612 100 2.607

2.345 2 2.373

2.259 2 2.280

2.042 100 2.064 100 2.071 2.006

1.840 10 1.847 50 1.818 1.841

1.503 3 1.509 20 1.5063

1.433 7 1.433 20 1.4336

1.222 4 1.291 10 1.2215

1.179 21 1.179 10 1.1703

1.163 2 1.167 7 1.168 20 1.164

1.063 1 1.065 7 1.066 16 1.051 1.061

1.021 7 1.034

0.914 2 171

-7 Specimen sulfidised at T = 1173 K. P = 10 atm. b2

Sample FeS oc - MnS FeAl2S4 MnAl2S4

I/I. I/I. I/I. d(A) (%) d (A) (%) d (A) (%) d(A) d (A)

1.056 9 1.048 10 1.0662 16 1.051 1.0614

1.103 6 1.101 10

1.156 14 1.167 7 1.1682 20 1.164

1.241 8 1.263 1.2215

1.296 9 1.291 10 1.306 8

1.324 10 1.321 20

1.391 9 1.3836

1.427 9 1.433 20 1.4336

1.445 10 1.442 10

1.500 27 1.490 10 1.509 20 1.487 1.5063

1.556 8 1.575 6 1.569

1.611 10 1.606 7 1.653 1.673

1.726 54 1.720 40 1.754 1.724

1.831 75 1.847 50 1.818 1.841

2.069 88 2.064 100 2.071 2.096

2.212 9 2.207 7 2.280

2.607 100 2.612 100 2.576 2.607

2.660 52 2.640 50 2.683

2.763 4 2.715

2.885 10 2.870 10 2.879

3.015 80 2.980 40 3.015 14 2.977 3.018

3.017 4 3.107 3.134 172

Specimen sulfidised at 1173 K, P atm. b2

Sample FeS oc - MnS d(A) I/I, (%) d (A) i/i, <%) d (A) I/I, (%)

1.050 12.5 1.048 10 1.0662 16

1.145 15.6 1.1682 20

1.325 6.3 1.321 20

1.442 5.0 1.442 10

1.486 6.3 1.490 10 1.509 20

1.599 12.5 1.606 7 1.575 5

1.724 20.0 1.720 40

1.824 62.5 1.847 50

2.074 40.6 2.064 100

2.574 100.0 2.612 100

2.640 25.0 2.640 50

2.972 18.8 2.980 40 173

Specimen sulfidised at T = 1173 K, PQ 10 atm. b2

Sample FeS oC - MnS MnAl?S4 FeAl?S4 I/I. I/I. I/I, o O O O o d (A) (%) d (A) (%) d (A) (%) d (A) d (A)

1.051 17.9 1.048 10 1.0662 10 1.0614 1.051

1.105 15.5 1.105 15 1.1682 20 1.1703

1.296 3.1 1.291 10 1.303 1.263

1.327 20.2 1.321 20 1.306 8

1.433 8.5 1.433 20 1.4336

1.445 8.5 1.442 10

1.493 3.1 1.490 10 1.509 20 1.5063 1.437

1.612 5.4 1.606 7

1.645 3.1 1.636 10 1.653

1.682 1.6 1.673

1.728 76.1 1.720 40 1.724 1.754

1.830 13.2 1.847 50 1.841 1.818

2.075 100.0 2.064 100 2.096 2.071

2.576 7.8 2.607 2.576

2.661 59.0 2.640 50 2.612 100 0 2.683

2.759 4.7 2.715

2.887 5.0 2.87 10 2.879

3.001 57.5 2.98 40 3.015 14 3.108 2.977 174

Specimen sulfidised at T = 1173 K, P atm. b2

Sample FeS oC - MnS FeAl2S4 MnAl2S4

I/In i/i> i/i. d(A) (%) d (A) <%) d(A) (%) d (A) d (A)

1.053 12 1.048 10 1.0662 16 1.051 1.0614

1.086 4 1.065 7 1.097

1.113 13 1.105 15

1.123 8 1.125 7

1.140 8 1.167 7

1.183 4 1.174 10 1.1682 20 1.1703

1.232 7 1.2215

1.281 4 1.291 10 1.263 1.274

1.333 22 1.321 20 1.306 8

1.393 4 1.433 20

1.455 15 1.442 10 1.460

1.506 9 1.490 10 1,509 20 1.487 1.506

1.631 4 1.636 10 1.634

1.692 5 1.653 1.673

1.730 55 1.720 40 1.724

1.758 6 1.769 7 1.754 1.769

1.818 3 1.818

1.842 11 1.847 50 1.841

1.945 14 1.987 7 2.006

2.095 100 2.064 100 2.071 2.096

2.176 8 2.207 7

2.338 6 2.373

2.557 5 2.576 175

------Sample FeS ex - MnS FeA^S^ MnAl2S4

I/I. I/I, I/I, d(A) (%) d(A) (%) d (A) (%) d (A) d (A)

2 690 69 2.640 50 2.612 100 2.683

2.797 4 2.870 10 2.879 2.715

3.015 62 2.980 40 3.015 14 2.977 3.018

3.175 6 3.107 3.171

3.463 5 3.440 7

3.982 3.98 4.012