jH-'xc - ^ £- \

MHD MATERIALS - SEED/SLAG INTERACTIONS AND EFFECTS

QUARTERLY PROGRESS REPORT

January - March 1980

Samuel J. Schneider

Project Manager

Center for Materials Science

U. S. Department of Commerce

National Bureau of Standards

Washington, D. C. 20234

PREPARED FOR THE UNITED STATES

DEPARTMENT OF ENERGY

MHD OFFICE

"This report was prepared as an account of work sponsored by the United

States Government. Neither the United States nor the United States

Department of Energy, nor any of their employees, nor any of their

contractors, subcontractors, or their employees, makes any warranty,

express or Implied, or assumes any legal liability or responsibility

for the accuracy, completeness, or usefulness of any information,

apparatus, product or process disclosed, or represents that its use would not infringe privately owned rights."

TABLE OF CONTENTS

Pa^e

I. SUMMARY OF PROGRESS TO DATE 1

II. DETAILED DESCRIPTION OF TECHNICAL PROGRESS 3

1. Thermochemistry of Seed and Slag 3

2. Electrical Conductivity and Polarization 18

3. Corrosion of Downstream MHD Components . 26

i

I. SUMMARY OF PROGRESS TO DATE

1. Thermochemistry of Seed and Slag

Vapor pressure data of K(g) over KA1S104 as a function of bulk

composition and 1/T are reported. The bulk composition range varies

from 28.9 to 10.7 wt% K2 0 while the K pressure decreases by 3 orders

of magnitude. Some non-equilibrium effects are shown for compositions

containing only small amounts of KA1 SI

0

4

. Mullite was the only

crystalline residue identified at the conclusion of the experiment.

Vapor pressure data of SiO(g) and 02 (g) over the KA1 Si

0

4 showed an activity of Si

0

2 close to unity. Measurements of the K pressure

over KA1 Si

0

4 and calcium silicate salts indicated complete removal

of the K2 0 at 1945 K.

Additional evidence supporting the existence of (KA102 -K2 C0 3 ) ss

is presented. Melts with known phase equilibrium properties in the

system K2 0-Ca0-Al 2 03 -Si02 are being prepared for vaporization

studies. A quasichemical melt polymerization model for treatment

of K2 0-Ca0-Fe0 -Al 2 0 3 -Si02 equilibria is discussed.

A

2. Electrical Conductivity and Polarization

The electrical conductivity of YCr03 doped with 5% Ca has been

measured in a highly reducing atmosphere at temperatures from 1300 to

1600 °C. The samples are not single phase In composition, however,

and this may have caused some errors in the bulk conductivity values

due to hair-line cracks or phase boundaries.

Several polarization experiments have been completed using slag

(Bow, NH) as an electrolyte and using platinum as electrodes. These

were made in a specially designed cell incorporating four probes to

allow polarization to be measured at any probe as a function of time

at a given temperature and for a preset condition regarding charge

transfer (constant current or constant voltage).

The electrical conductivity of a Bow, NH slag sample containing an added amount of iron has been measured. At lower temperatures (1200 °C)

the conductivity increases nearly proportional to the added iron but at

higher temperatures (above 1350 °C) the conductivity increased by a factor of 5 when the iron content was doubled. This increase may be

due to the decrease in viscosity associated with iron addition and/or

equivalent silica decrease.

3. Corrosion of Downstream Components

Type 316 stainless steel tubes, cooled to 400 °C, 500 °C, and 590 °C,

were exposed to oxygen rich hot gas streams seeded with K2 S 04 . Optical,

SEM, and EDX analysis indicate the geometric and chemical nature of the

deposits and reaction zones to be similar to those of Type 304 stainless

steel exposed previously under the same conditions. As in those cases,

Fe appears to have penetrated the farthest into the salt deposit region

while the dispersion of Ni in the reaction zone area remains the most

difficult to follow. The extent of wastage (weight loss) of Type 316 and

Type 304 stainless steel exposed at 590 °C and 400 °C to K2 S04 seeded

oxygen rich hot gas streams was also determined.

1
In the coming quarter further exposure tests will be conducted on Type 304 and Type 316 stainless steels to determine wastage under oxygen

rich conditions with a seed of 80 percent by weight K2 C04 and 20 percent

by weight K2 SO4

.

2
II. DETAILED DESCRIPTION OF TECHNICAL PROGRESS

1. Thermochemistry of Seed and Slag (E. R. Plante and L. P. Cook)

A. Vaporization Studies (E. R. Plante)

Progress : Vapor pressure measurements on KA1 Si

O

4 were concluded.

Measurements on a selected composition in the K2 0-Ca0-Al 2 0 3 -Si02 system were started.

The K2 0-A1 2 0 3 -Si02 system has the potential for forming several

stable or metastable phase assemblages in which the K2 0 activity is

fixed according to the phase rule. These include KAlSi 04 , K2 0*9A1 2 03 , KAlSi 2 0gj A1 2 0 3 j 3A1 2 0 3

• 2S i

0

2 ; and KAlSi 2 0g, KAlSi 3 0g, 3A1 2 03

• 2Si

While analysis of the experimental results obtained is not entirely

complete, no evidence that the K pressure was independent of bulk

composition of the sample was obtained. This is in agreement with previous

studies from this laboratory in which constant potassium pressures were

not observed even though KAlSiOi+, K2 0*9A1 2 03 and KA1 Si 2 06 crystalline

phases were identified as being present in the sample. This represents non- equilibrium behavior in the measurements. Other evidence for non- equilibrium behavior is shown for the experiment in Figure 1.

Here, the observed K pressures are greater when the temperature is

decreased even though the bulk K2 0 composition decreased in the first

part of the experiment. This effect is likely due to depletion of

KA1 Si

O

4 from the sample. At the highest temperature (1880 K) partial

melting probably caused a redistribution of some remaining KA1 Si O4

by either bringing the crystals to the surface or dissolving them in

the melt. This effect was noted only in the experiments for which

the remaining KAlSi 04 phase was becoming small.

Some representative vapor pressures of K as a function of 1/T

and bulk K2 0 concentration are given in this report. Figure 2

shows K pressure data as a function of 1/T at various bulk K2 0 com- positions. The data shown was obtained during experiments in which

the sample temperature was decreased throughout the run. Bulk K2 0

concentrations were calculated from the integrated ion current data

(1 At/T) and represent the K2 0 concentration that would exist if the

sample were a homogeneous solution. The calculation of bulk com- position did not include loss of Si

0

2 which became appreciable toward

the end of the experiment at higher temperatures. Linearity of the

vapor pressure data with 1/T was better for the experiments in which

the temperature decreased than for those in which the temperature was increased. This may be due to better mixing of the KAlSi 04 on melting.

The data shown represent an extremely wide range of concentration and

temperature. The K pressure at constant temperature decreases by over

3
3 orders of magnitude when the K2 0 bulk composition decreases from

29 wt% K2 0 to 11 wt% K2 0. At the highest temperature, 2160 K, the

sample should be entirely molten and under these conditions slow

kinetics due to complex crystal chemistry is not a possibility. The

activity of K2 0 in the melt is so low that it leads to the question

of whether the K2 0 concentration could be as high as 11 wt%. Assuming

an error of 30% in the mass spectrometer constant, the K2 0 concentration

might be as low as 3 wt%.

During a considerable portion of the experiment, measurements of

th| SiO intensity were possible and these data along with measured

02 intensities were used to calculate the product P 5io'F> 0'' * W^ 1C ^

is related to the activity of Si

0

2 and K for pure Si0 2 U). These

data showed that the activity of Si

0

2 is p close to unity for this

system.

The experiment was finally terminated because the P<»

.

q

was greater

than and it appeared that the sample might become exhausted in Si

0

2 before all of the K2 0 was evaporated. The sample consisted of a very

hard crystalline mass embedded in a glassy matrix. Heating in hot

(100°C) dilute HF (^30% cone. HF in H 2 0) dissolved the glassy matrix

but not the crystalline mass. An x-ray diffraction pattern showed

the crystalline material to be mullite (3Al 2 0 3 *2Si02

). It is presumed

that in the final stages of the experiment, crystalline mullite was in

equilibrium with a liquid phase containing K2 0, Si

0

2 and A1 2 0 3

.

Measurements on a sample containing 49.84 wt% KAlSi04 , 16.47%

Ca 2 Si04 , 1 6.47% Ca 2 Al 2 Si07 and 17.21% Ca 3 Si 2 0 7 were begun. The

sample contained each of the above phases in an intimately mixed form.

The early stage of the evaporation is expected to be described by the

reaction

2 KA1 Si 04 (c ) + 7Ca 2 Si04 = 4Ca 3 Si 2 0 7 + Ca 2 Al 2 Si07 + 2K(g) + 02 (g) (1

)

Figure 3 shows K pressure data during the early stages of evaporation

up to the point where the Ca 2 Si04 phase should be exhausted. In addition, pressures calculated from the literature at 1400, 1600 and

1800 K are shown for comparison. The agreement between the experimental

and calculated pressures are reasonable although the experimental

slope is significantly steeper. Additional measurements on this com- position were made. It proved possible to vaporize all the K2 0 from

the sample at 1945 K. This appears to confirm the suspected effect

of CaO in ameliorating K2 0 capture by A1 2 0 3 and Si

0

2

.

4
B. Phase Equilibrium Experiments (L. P. Cook)

1 . System K20-Ca 0-Al 2 0 3 ” S i

O

2

For vaporization studies, a composition consisting of 49.8 wtX KA1S104 , 16.5 wt% Ca 2 Si04 , 17.2 wt% Ca 3 Si 207 , and 16.5 wt% Ca 2 Al 2 Si 07 was prepared.

It was thought that this would give a large range of possible potassium

pressures [1]. Measurements will be discussed under vaporization studies

in the next report.

For additional vaporization studies, it is desirable to synthesize

mixtures with a well-defined melting point. It is necessary to use invariant melts for this purpose. The approximate composition of

several such melts are given in Table 1. Compositions of these and

other invariant melts will be reported with improved accuracy, pending

full reduction of EDX data.

Table 1. Approximate Compositions of Selected Invariant Melts,

System K

2

0-Ca0-Al ^O^-Si

C

2

(wt %)

.

Solids Present

at Melting CaO ai

2 o3

Si0

2

KAlSi0

4

/CaSi0

3

11.7 25.7 19.2 43.5

CagAl

2

Si0

7

/CaSi0

3/

CaAl

2

Si

2

0

8/KAlSi2

0

6

6.0 30.3 25.4 38.4

KAlSi0

4

/KAlSi

2

0

6

/

Ca^l

2 SiOy/CaSiC

3

10.5 25.5 24.0 40.0

Ca^SiCyKAlSiCy

KAlSi

2

0

6

/CaAl

2

Si

2

0

g

11.5 17.6 33.7 37.2

Ca^^SiOy/Ca^SigOy/

CaSi0

3

/KAlSi0

4

6.5 40.1 17.1 36.4

KA!Si0

4/Ca

2

Al

2

SiC

7

11.8 23.6 37.4 27.2

5
2 . System K2O-C3O- FeO^-Al 203-Si O2

The effect of a slag coating on the electrical performance of

the generator in open cycle coal-fired MHD systems continues to receive

a substantial amount of attention. The problems associated with the

behavior of slag in downstream components, and with the recovery of

potassium seed captured by slag areno less significant in the operation

and design of economical MHD systems. A quantitative knowledge of

slag properties appears to be essential for several aspects of MHD

research and development.

In an earlier paper

1 the available data on MHD channel slag

compositions were summarized and working models were proposed for the

bulk composition of slags produced by MHD combustion of typical subbituminous ("Western") and bituminous ("Eastern") coals. Vapor pressure

measurements were reported for these two synthetic slags. The present

paper is a continuation of this work, and here are described the results

of exploratory phase equilibrium studies on the model slags. Also

presented is an attempt to make meaningful generalizations and extrapolations through theoretical modeling of this experimental data. It

is hoped that ultimately the otherwise intractable complexities inherent

in the physical chemistry of MHD slag will become better understood so that reliable prediction and control of viscosity, electrical and

ionic conductivity, seed retention, melting and crystal 1 iztion behavior,

corrosiveness, and other compositionally dependent properties will be

possible.

In preparation for the development of a predictive model for immiscible

K« 0 -Ca 0 -FeC

x -Al

? 0 ^-Si Cp liquids, theoretical manipulations were begun

or data takin previously (ref. [1]) on compositions of melts coexisting

with KAlSiC^ [ 2 ] for bulk compositions modeling "Eastern" and "Western"

channel slag. It is expected that these results will be presented at

the 7 th International MHD Conference. The following paragraphs are excerpts from a paper submitted to that conference.

The ideal solution provides a useful starting point for any

modeling effort.. One form of the ideal solution equation for binary

crystal /I i quid equilibrium in system A-B is:

, fusion

-In XTM

elt = -In X*

1 (} - \

fusion)

B

where xTu and Xg

1 are tbe mole fractions of component B in coexisting

liquid and crystal, AHg

USlon

is the enthalpy of fusion of pure crystalline

B, Tg

US on

is the melting point, R is the gas constant, and T is the

temperature in degrees kelvin. It can be shown that this relation

applies as well to crystal /I i quid equilibria in multicomponent systems,

In performing tTiis calculation for the "Western" melts of ref. [ 1 ],

the relatively small amount of iron in the residual melts may be assumed

to be in the ferrous state. Sample calculations indicate that for

purposes of this paper the above assumption is not likely to result in

significant changes in the results of theoretical manipulations or their

interpretation (for the "Eastern" melts this is emphatically not true).

6
The ideal solution calculations include the 85 mol % purity of the

crystals involved, the heat of fusion used by Spencer et al . [3],

and an assumed congruent melting point of 2058 K. The resulting curve is

shown in Figure 3a. If the precipitation of KAlSiCL is assumed to be

defined by

(KAlSi0

4

) x1 = l/2(K

2 0)

melt

+ l/2(Al

2 0

3

) melt

+ (Si0

2

) melt

so that

j^ea!, me lt . melt 1/2 , Ymelt xl/2 / Y meltx

a KAlSi0

4

" SlA

K

2

0

; Al

2

0

3

j lX Si0

2

j

then the calculated activities in Figure 3a are produced from compositional

data in ref. [1]. The agreement between predicted and measured activities

is thus not good, if the observed activity is defined in this way.

Some success in explaining the non-ideality of silicate melts has

resulted from application of polymer theory. According to the model

proposed by Toop and Samis [4], polymerization reactions may be sum- marized in quasi chemical form by:

-

0 20 =0 +0

2 - -

0 where o , 0 and 0 represent free oxide ions, non-bridging oxygens

and bridging oxygens respectively. An equilibrium constant, k, can be

defined accordingly. Values of k from the literature [4,5], with a temperature dependence assumed similar to that of the Fe0-Si0

2 system

[6] have been combined in the present work to calculate the distribution

of quasichemical oxygen species in the melts of ref. [1]. A correction

for sUnda restate differences, in proportion to the ion field strengths

of Ca Fe , ^nd Mg

L has been applied [7]. Otherwise values of K

for Fe

L and Mge‘ are assumed equal. In these and subsequent calculations

KA10

2 and Si02 are considered as substituting ideally in the polymer

network [8] e ^By solving the following equations, the number of moles,

n, of 0", 0 , and 0

d ~ per mole of melt may be found, where X is the

mole fraction of ( Si

0

2 + KA10

2 ):

(8X - 8X

2 )n

2 + (-2X -2)n + (1 - 4k) = 0

O' O'

4X - (n _)

n o = 0

n

? = (1 - X) - ( O

0 2

Results of this series of calculations for the melts of ref. [1] show

marked increase in polymerization with increasing (KA10

2 + Si

0

2

) , as indicated by the increased 0 content of the melt. The free oxide

ion content is, as would be expected, very low for all compositions.

7
Fraser [7] has modified the theory of Huggins [9] to allow prediction of the proportions of five basic Si

0^

structural units

present in the melt using the quasichemical species of the Toop-Samis

model. Structural units are indicated schematically by numbers in

parentheses which give the number of non-bridging oxygens. Thus

Si0.(0) represents a structural unit of the infinite framework, while

Si0?(4) is the independent Si 0^ tetrahedron. Results of calculations

for^the residual melts show that only the Si0.(0) and SiO^(l) units

increase in abundance with increasing mole % (KA10

2 + SiCL), and hence,

with increasing temperature. If an ideal activity is defined according

to

ideal, melt =

. / Y nelt % , Y (0), meltv

a KAl Si 0

4

UKA10

2

; A Si0

2

'

where is the mole fraction of (KA10

2 + Si

2

) in the melt

present in^the form of Si0 4 (0) structural units, or structons [9],

then the calculated points in Figure 6b result, giving substantially

lower activities than those predicted. However, if the calculation,

is completed according to

ideal, melt . / Ymelt v , Y (1), meltv

a KAlSi0

4

'

4 ' a KA 10

2

1 Si0

2

‘

then the calculated KAlSiCU activities for the "Western" melts agree very well

with those predicted from ideal solution theory, as shown in Figure 3c.

It is suggested that while Si0 4 (0) is the structural unit most

similar to the KAlSiC 4 structure, Si0 4 (l) has a non-bridging oxygen

essential for reaction and therefore is the melt structural unit

involved in the precipitation equilibrium. The result of this

calculation may be taken as evidence that KA10 2 and Si0 2 do indeed

mix in a near ideal fashion in the polymer network; that KA10 2 in

all sites in the polymer network is of nearly equal reactivity; but

that it is the degree of polymerization of the silica within the

polymer network which determines the precipitation reaction of KAlSiOi*.

The quasi chemical polymerization model can also be used to calculate

the free energy of mixing for "Western" melts (Figure 3d) according

to the following equation:

AG

mi xi ng

n 0~ RT In k

2

8
support of the stability of these

Note, however that the

A nearly flat curve results in

melts with regard to unmixing.

temperature data show irregularities between 1350 and 1550°C.

This may be an indication of the misicibility gap known to occur

a short distance away in compositional space, at the site of the

"Eastern" bulk composition.

vs

The approach has also been extended to results of experiments

on the "Eastern" iron-rich bulk composition of ref. [1]; hov/ever,

the liquidus surfaces are apparently too complex to be modeled by simple

precipitation reactions of the type outlined above, which worked

reasonably well for the iron-poor, calcium-rich "Western" composition.

This probably is the result of non-ideal entry of ferric iron into the

silicate polymeric framework. Nonetheless the quasi chemical approach

seems to work well for description of the free energy of mixing, where

it is possible to ignore specific polymeric interactions. To perform free

energy calculations on "Eastern" melts, a series of compositions along a vector through seven-component space have been chosen. The vector

is defined by the two immiscible melts coexisting at 1550°C (ref. [1])

whose compositions are given in Table 2. Ferric iron and A1 2 0 3 have

been included in the calculations by assuming equilibrium constants

appropriate for relatively undisscciated oxides (3.81 and 4.38,

respectively). Results of these free energy calculations are shown

in Figure 7, at a series of Fe+ 3 /Fetot£l ratios. It has been assumed

for purposes of these calculations that this ratio is the same in

both melts of the immiscible pair. Figure 7 illustrates the rather

dramatic effect of Fe+3 /Fetotal on the mixing properties of these iron

melts. A large region of liquid unmixing for more oxidized melts is

predicted. Composition of predicted and observed melts can be compared

in Figure 7b, and it can be seen that by judicious choice of equilibrium

constants and assumed Fe+3 /Fe tcta ' reasonable agreement can achieved

through use of the quasichemical model.

The quasichemical approach offers hope in the development of a quantitative model for description of slag behavior under MHD conditions

Further progress will require knowledge of how entry of ferric iron

into the melt affects polymerization reactions.

9
. Compositions of Coexisting Immiscible

"Eastern" Bulk Composition (mole %).

Liquids at 1550 °C,

Silica-rich (S) Iron- rich (F)

CM W> O 45.87 25.07

A1 2°3 15.58 9.00

Fe0

x 18.57 49.85

CaO 4.92 6.88

MgO 1.30 1.32

k

2 o 13.76 7.88

100.00 100.00

Plans Continue vapor pressure measurements on selected portions of

the K2 0-A1 2 0 3 -Ca0-Si02 system.

Preparation of melts with known phase equilibrium behavior for

vaporization studies in the system K2 0-Ca0-Al 2 0 3 -Si02 will continue.

Studies of subsolidus equilibria in the system will resume, and will

concentrate upon equilibria involving KAlSi04 , Ca 2 Si04 , K2 CaSi04

( KA1

0

2 -K2 C0 3 ) ss and related Dhases.

References

1. W. R. Hosier, T. Negas and S. J. Schneider- Development- Testing and

Evaluation of MHD Materials, NBS reports prepared under DoE Contract

EA-77-A-01-6010, Aoril-June 1978 and July-September 1978.

2. L. P. Cook, R. S. Roth. H. S. Parker, and T. Negas, Am. Mineral.,

62 1180 (1 977).

3. F. E. Spencer, J. C. Hendrie. dr., and D. Bienst.ock, Sixth Int.

Conf. on MHD Electrical Power Generation, 2 181 (1975).

4. G. W. Toop and C. S. Samis. Trans. Metal. Soc. AIME, 224 878 (1962).

5. C. K. Masson, I. B. Smith, and S. G. Whitewav, Canad. J. Chem..

48 1456 (1970).

6. P. A. Distin, S. G. Whiteway, and C. R. Masson, Canad. Metal.

Quart. ,. 1_0 73 (1 971 )

.

7. D. G. Fraser, p. 301-325 in: D. G. Fraser, (ed)., Proc. NATO

Adv. Study Inst., Oxford, England, 1976 (1977).

8. F. J. Ryerson and P. C. Hess, Geochim. et Cosmochim. Acta., 42_

921 (1978).

9. M. L. Huggins, J. Phys. Chem., ^58 1141 (1954).

10.

E. R. Plante and L. P. Cook, Proc. 17th Symp. Eng. Aspects MHD, Stanford,

C.1.1. (1978).

10
Potassium

Pressure

(atm)

Figure 1. Non-equilibrium effect in the vaporization of KAlSiOir due to

low concentration of KAlSi0 4. and partial melting.

11
Potassium

Pressure

(atm)

Figure 2. Selected K vapor pressure data of KAlSiOu as a function of 1/T

and bulk composition.

12
Potassium

Pressure

(atm)

Figure 3. Vapor pressure of K produced by reaction of KAlSiOi,

with Ca 2 SiOi*.

13
r

n K 2°"

Figure 4 . Schematic phase diagram showing stability of (KAIO2-K2CO3) ss.

14
20 (CuKa)

Figure 5. Low angle portion of powder pattern of (KAIO2 -K2 CO 3 ) ss

15
Activity

of

KAISI0

4

Activity

of

KAISIO,

T(°C) T (°C)

.2600 r

.2200

.1800

.1400

.1000 -

.0600

.0200

Ideal Solution —

/

//'

/

/

*

4(X I (X m'" i ^

"'‘kAIOj1 1 SiOj > y

/

* *

1300 1350 1400 1450 1500 1550 1600 T(°C) Mole % (KAI02 + Si02

)

Figure 6. Theoretical manipulations of melt data in ref. 1, "Western"

model slag, a) Ideal solution freezing point depression

compared with activity of KAlSiOi* calculated from precipitation of the oxides, b) Comparison of ideal KAlSi0 4 activity with that calculated for precipitation using Si0 4 (0)

structons. c) Comparison of ideal activity with that predicted from precipitation by Si0 4 (l) structons. d) Calculated

free energy of mixing for "Western" residual melts as a function of mole % (KA102 + Si0 2 ).

16
mixing

x(KAIO, + Si02)

X(KAIOj + Si02

)

x(KAI02 + SiOj)

b

d

f

Figure 7. Calculated free energy of mixing along a compositional vector

passing through the compositions of immiscible melts, "Eastern"

bulk composition (see Table 2). Calculations assume 1550°C

and Fe+3 /Fetotal = 0.3000 (Fig. 7a), 0.2500 (Fig. 7b), 0.2000

(Fig. 7c), 0.1500 (Fig. 7d), 0.0909 (Fig. 7e) and 0.0476 (Fig.

7f). Tangency is drawn in Fig. 7b to facilitate comparison

with observed immiscibility (melts S and F). These free energy

curves have been drawn by computer to fit calculated points

using a piecewise cubic interpolating polynomial. Dashed

portions of the curves lie slightly off the compositional

vector.

17
2. Electrical Conductivity and Polarization (W. R. Hosier)

Progre ss

a. Electrical Conductivity of YCr 03

Initial measurements of the pressure dependence of the electrical

conductivity of YCr 03 doped with 5% calcium on the yttrium site have

been completed. There has been some difficulty in establishing the

oxygen partial pressure in the sample chamber that corresponds to the

oxygen partial pressure in the flowing gas where oxygen/nitrogen gas

mixtures are used. This difficulty arises because of the large sample

chamber (1200 cm 3

) originally filled with air which must be swept out

with the flowing environment gas of chosen composition. True oxygen

pressure dependence of the conductivity can only be measured after the

environment gas has been completely changed in the sample chamber. This

requires times in excess of 4 hours with a flow rate of 450 cm 3 /minute.

In contrast, however, for those atmospheres where very low oxygen pressures

are obtained by the reaction 2H 2 + 0 2-* 2H 2 0 such as is obtained using

forming gas (95 nitrogen to 5 hydrogen), the low oxygen pressures are obtained quickly (1 hour or less) in the environmental chamber. The

measurements have been completed in atmospheres containing about 5 ppm

02 in N 2 and in an atmosphere of 10~ 25 atmosphere oxygen as indicated on

a zirconia tube oxygen sensor gauge based on the Nernst-Einstein relation.

(Intermediate oxygen pressures between 5 x 10~ 6 and 10" 25 atmospheres

will be obtained in future experiments by mixing the 5 ppm gas with

forming gas.)

Table 1 gives values of conductivity and soak temperatures. The

conductivity values are those taken at the end of the soak period, but

data were taken at 5 to 10 minute intervals during the soak period, before

equilibrium was established. The order in the Table is the sequence

in which the data were taken. As can be seen, the scatter in the conductivity

data is quite large.

Temp. °C

0

2 pressure

^(atmos.

)

Table 1

Conductivity

a (ohm-cm)"

Total flow time

before meas. (hr)

Soak Tim

at T. (h

1300 5 x 10' 6 2.9 18 18

1500 II 3.12 20 2.5

1400 II 3.05 23.5 3.0

1400 5 x 10" 25 1.05 11.5 11.5

1600 II 1.30 14.5 3.0

1500 II 1.05 16 1.5

1400 II 0.92 17 1.0

1300 II 1.04 36.5 19.5

1600 II 1.02 45.0 8.5

1500 II 0.88 56.5 11.5

18
There has been some difficulty In preparing good single-phase

material. Possibly, as a result of this multi-phase condition, samples

oxidized after having been prepared in a highly reduced state sometimes

exhibit small hairline cracks. These cracks will cause erratic results

in the conductivity values.

YCr0 3 , doped with Ca or Mg in a highly reduced state, is very hard

and it is difficult to saw or drill with diamond tools. In the oxidized

state, however, it is relatively soft and sometimes fragile and breaks

easily upon handling. More work must be done on phase composition and

microstructure and their relation to oxidation/reduction conditions

before final work can continue on the electrical conductivity measurements.

b. Electrochemical Effects of Slag-Electrode Interfaces

An electrochemical cell, consisting of four platinum probes immersed

in Bow, NH slag containing 20% added K2 SO4 , was constructed. This cell was heated to 1250 °C and subjected to a constant applied voltage mode

test where the voltage was reversed periodically during the course of a 2-1/2 hour run. Voltages were monitored on all probes during the test.

Figure 1 is a radiograph of the electrochemical cell, while Figure

2 shows the voltage on the various probe sets as a function of time. A

constant 31 volts was maintained between the top and bottom probes and

this resulted in an initial current, after each current reversal, of 60 to

100 mA, which drops off with time as shown in the figure. The current

was reversed at 5 minute intervals (after a initial interval of 2 1/2

minutes) over a period of 2 1/2 hours. Only 23 1/2 minutes of the data are shown, but the remaining data over the entire experiment period are repetitive. It can be readily seen that, for example, when the top

probe is positive (anode), as in periods A, C, and E, nearly all the potential

drop in the cell is across the Iy p t

set of probes, after a few minutes

operation, while the voltage across Ig pn (In cathode) is low. Current

reversal then produces a large potential drop at the I„ „ probes and a small voltage at the I T T probes. In previous experiments (see QR

Sept. -Dec. 1977), some iWn depletion was observed in the anode area just

outside the platinum electrode in an electrochemical cell run for a period

of 20 hours at 1400 °C. At the same time, definite enrichmen^ of iron was observed at the cathode slag-electrode interface. Cadoff , in

electrochemical work on slag-electrode reactions, shows that in cells

designed for static slag (non-flowing), the potential buildup of the cell

Is due to the accumulation/of electrically resistive ions at the slag/anode

interface. Koester et al . ' has shown that the critical potential for

arcing or breakdown is higher for AISI 1018 steel than for nickel or Type

304 stainless steel. This may be due to the supply of iron from the anode

electrode alleviating the usual iron poor (or silica rich) area just at

the anode/slag boundary, usually associated with an electrode material

not containing iron. This experiment demonstrates that this potential

buildup is relatively rapid and is associated with the slag-anode interface.

At the same time, the conductivity of the bulk slag remains unchanged

since the potential on the probes pj p

B drops off in proportion to the

drop off in the primary current while the applied voltage is constant.

19
A second experiment on this same cell was performed using a constant

current (30 mA) mode but reversing the current periodically (5 minute

intervals). Figure 3 shows that data obtained over a 20 minute period

in the experiment. Although the experiment was run for 1 1/2 hours,

only slight changes were observed with time on reversal of the current.

Breakdown in the current conducting mechanisms, after a current reversal,

appears to occur earlier, when the I B probe is an anode as in period B and D (1.5 to 2.0 minutes), than when the I T probe is an anode as in

periods A and C (4 to 5 minutes). Since thi experiment was run at

1250 °C, there could be a difference in the oxidation state of the iron

between the top and bottom of the cell and this might cause a difference

in the breakdown potential when the anode is at the bottom of the cell

(away from ambient atmospheres). Also, with the top of the cell as anode

for a 5 minute period followed by a 5 minute period with the current off

(these periods not shown in figure), the cell returns to its original

state;, i.e., the potentials are distributed evenly along the length of

the cell, showing that any depletion layer has been repaired. However,

with the bottom of the cell as anode, and the same procedure followed,

the depletion layer is not repaired completely even after a period of 50

minutes. A measurement of all probe voltages with current off yields

zero potential indicating that there is no measurable battery effect

present. In the constant current mode, the high potential develops

between the anode and its adjacent conductivity probe, as was also observed

in the constant voltage mode. Breakdown occurs also at the anode. The

corresponding cathode potential decreases with time after a slight

initial increase after reversal. The slag bulk conductivity decreased

(potential on p probes increased) slightly during the course of the

experiment which may be due to iron diffusion into the platinum electrodes.

Both of these experiments were done at a temperature (1250 °C) where

the slag is very viscous and where there may be a combination of ionic

and electronic (semiconductive) conductivity. If time permits, this

experiment should be repeated at higher temperatures where the conductivity is believed to be largely ionic and due primarily to the movement

of iron ions. '

c. Slag Electrical Conductivity

A sample of slag was made from the Bow, NH slag which contained an added 20% K2 SO4 , but now with an additional 20% of Fe 2 0 3

. Figure 4

shows the conductivity data as a function of temperature in an air

atmosphere. The sample was heated to 1500 °C and held at that temperature

until the conductivity became constant as a function of time (< 1/2

hour). It was then cooled as rapidly as possible in the furnace to

900 °C which required a period of 24 minutes. The temperature was then

lowered slowly to 500 °C. The data shown in Figure 4 were taken with

temperature increasing; conductivity versus time was monitored to obtain

maximum or minimum conductivity values. The conductivity is higher in this

sample than in its companion sample without the added iron, shown in

QR July - Sept., 1979. The conductivity near the maximum at 1200 °C is

approximately a factor of 2 higher in this sample with the added iron

when long cool down times through this temperature range allow for crystallization of the slag. X-ray diffraction measurements indicate that at

20
these conductivity maxima, the crystalline components of the slag are feldspar sesquioxide, and spinel solid solutions of undefined bulk

compositions. It is not possible to deduce whether the conductivity

is due to these crystalline phases or to an iron-rich glassy matrix

likely to be surrounding the crystalline components. In the latter case,

the conductivity would be predominantly ionic. However, polarization

effects were not observed in this temperature range during the electrical

measurement.

The conductivity in the higher temperature range above 1350 °C is a factor of 5 greater in this sample with added iron >,(24.2 w/o) than in

the sample with no added iron (14.7 w/o). Capps^

4

' has shown that there

is a decrease in slag viscosity with increased iron content and with

decreased silica content. The relative silica content necessarily

decreases with any added component other than silica itself. This

disproportionate increase in conduci tivity may be dujSvto the viscosity

decrease reflecting the validity of the Walden rule'

3

' i.e. n • a= const,

where n is viscosity and o is the electrical conductivity for conducting

ionic species (iron ions) in a molten glass.

The data reported in this work are repeatable with repeated treatments and form the basis of a systematic approach toward understanding

the electrical conductivity in the coal slag which will coat the electrodes in an MHD generator.

PI ans

The electrochemical experiments will continue with the experiments

described in the report repeated at higher temperatures in order to

determine the polarization and relaxation times associated with any

depletion layer formed while the cell is carrying current.

Work on the slag electrical conductivity will continue also.

Preliminary investigations indicate that adding iron as Fe 2 0 3 or Fe 3 04 to Bow, NH slag without potassium causes excessive frothing and that

these additions can only be made in a highly reducing atmosphere.

Additional transition metal dopings may be investigated in order to

better understand the conductivity.

References

(1) Westinghouse Electric Corporation - Quarterly Report for the Period

Oct. 1 to Dec. 31 , 1977, Contract No. EX-76-C-01 -2248.

(2) Electrical Behavior of Slag Coatings in Coal Fired MHD Concentrations,

J. K. Koester, R. M. Nelson, 16th Symposium on Engineering Aspects

of Magnetohydrodynamics, University of Pittsburg, Pittsburg, PA,

May 16-18, 1977.

(3) Conductivity Mechanisms in Iron-Containing Slag, W. Hosier, G. White,

and T. Negas - 7th International Symposium on MHD, M.I.T., Cambridge,

Mass., June 16-20, 1980.

(4) Coal Slag Properties Related to MHD, W. Capps, Conference on High

Temperature Sciences Related to Open-Cycle Coal Fired MHD Systems,

Argonne National Laboratory, April 4-6, 1977.

(5) International Review of Science - El ectrochemistry- Physical Chemistry

Series Two, Volume 6, J. Bocknis, Editor, p. 290.

21
Figure 1. Radiograph of the electrochemical cell

showing the arrangement of the platinum

electrodes

.

22
Figure 2. Probe set voltages and current as a function of time with periodic applied

voltage polarity reversal. The applied voltage magnitude was constant

during the run. Cell temperature 1250 °C.

23
PROBE

VOLTAGE

(VOLTS)

Figure 3. Probe set voltages as a function of time with periodic applied

voltage polarity reversal. The current magnitude was constant

during the run. During breakdown, recorded voltage magnitudes

do not represent systematic cell behavior. Period A and C, was anode. Period B and D, 1^ was anode. Cell temperature 1250 °C.

24
a

)

CD

o

(D

h”

25

Figure

4.

Electrical

conductivity

of

Bow,

NH

slag

with

an

added

20

w/o K

2

S0

4

and

with

34.6

w/o

Fe20

3

(24.3

w/o

Fe)

as

a

function

of

temperature.

These

measurements

were

done

in

air.
3. Corrosion of Downstream MHD Components (J. Smit and C. D. Olson)

Progress : Tubular specimens of Type 316 and Type 304 stainless steel were exposed, under varied conditions, to K2 S04 -seeded , oxygen rich,

oxygen-propane hot gas streams. Following exposure, the specimens

were analyzed using SEM, EDX and in some cases gravimetric techniques

for indications of both incipient and gross corrosion. Test exposure

time, excluding apparatus warm-up and cool down time, was 4 hours.

Seeding with 250 grams of K2 S04 was effected during the first 25

minutes of that interval. Six stainless steel specimens were exposed

under these general conditions with the tube temperature maintained

at one of three levels: 400 °C, 500 °C or 590 °C. These temperatures

were fixed by internal air cooling and were monitored by a Pt/Ptl0%Rh

thermocouple welded into the specimen wall at the mid length point.

The specimens were seated in the test rig with the thermocouple

located in the center of the hot zone and facing the gas stream.

The gas stream temperature in the vicinity of the test specimen was

held at approximately 1300 °C as indicated by a Pt/Ptl0«Rh thermocouple.

A. SEM/EDX Analysis of the Metal -Deposit Interface

Metal lographic specimens were prepared from three of the Type 316

stainless steel specimens, 1-6, 2-6, and 3-6, exposed at 590 °C, 500 °C

and 400 °C, respectively. These sections were taken approximately

10 mm from the mid point position in the hot zone. To reduce the

possibility of external contamination and hydration in the coatings

formed on the tubular specimens during seeding, the specimens were

removed from the test rig upon cool down to about 40 °C and encapsulated

immediately in epoxy. The metallographic sections taken from the

encapsulated tubular specimens were cut, ground and polished using

non-aqueous media and were stored in evacuated desiccators prior to

SEM/EDX analysis.

On all the specimens a compact, somewhat columnar deposit formed on the

surface facing the gas stream. The thickness of this deposit varied

slightly with tube temperature and is approximately 1.7 mm. A much

thinner almost undetectable deposit, probably from fume, formed on the

trailing surface. These deposits are in the 0.2 mm range. At the

boundary, where the surface deposit and the fume deposit intersected

a thicker drip zone formed; see Fig. 1, 2 and 3. The SEM micrograph,

Fig. 4, shows a K2 S04 deposit on the leading surface of a Type 316

stainless steel tube held at 590 °C. This type of deposit is typical

of that produced from K2 S04 seeding and was seen on specimens of Type

304 stainless steel tested previously under these conditions. EDX analysis

indicates that potassium and sulfur is uniformly distributed through

the bulk of the deposit. An SEM micrograph, Fig. 5, shows the top of the

salt deposit with a fine whitish band on the surface. This band contains

a higher proportion of potassium than the bulk of the deposit possibly

due to vaporization of some sulfur at the deposit-gas stream interface.

26
The stainless steel deposit interface is seen in SEM micrograph, Fig. 6.

Note that the adhesion is rather poor as epoxy fills a gap in the

interface region, though this occurs after cooldown. A reaction area at the interface is shown in SEM micrograph, Fig. 7. At the metal edge, areas of low chromium and low nickel concentration are found dotted wit r regions of high nickel concentration, Fig. 7a - 7d. A loss of iron is

also associated with these areas. At the deposit edge, concentrations

of iron, chromium, potassium and sulfur are found indicating some degradation of the metal surface and formation of a reaction zone between

the metal and the deposit, Fig. 7e - 7h. The SEM/EDX analysis seems to

indicate a preferential leaching of the metal cations from the bulk material (Type 316 stainless steel) in a somewhat orderly fashion as in the case of Type 304 stainless steel. However, the extent of corrosion is

not as great as in the case of Type 304 stainless steel. This is

corroborated by wastage (weight loss) experiments also described in

this report.

As found in previous specimens, though nickel loss is readily detected,

its displacement to other areas is difficult to follow. Few areas are found where nickel has migrated even though its concentration is low.

Fig. 8 shows a reaction area where low chromium and low nickel con- centrations are found at the metal edge, Fig. 8a and 8b, and retal

cations are found at the edge of the deposit, Fig. 8c - 8e.

Migration of the iron, Fig. 9 , is farther back in the deposit and

generally parallels previous results in which iron was preferentially

leached out of the bulk stainless steel and into the deposit.

Fig. 10 shows a portion in the fume deposit area where the corrosion

process seems to be evident. On the metal edge low nickel areas are found, Fig. 10a - 10c, while areas of chromium, iron, potassium and

sulfur are found in the reaction zone. Again, nickel is the elusive

element. Though it is found in low concentration in the reaction zone,

it is not really found to have migrated into the deposit region.

At a lower temperature, 500 °C, corrosion is still observed in Type 316

stainless steel. Fig. 11 shows a region of corrosion at the metaldeposit interface. EDX analyses of the region, Fig. 11a - llh,

indicate areas of low metal cation concentration at the ~etal edge

and metal cation migration into the deposit. The same can be said

at the still lower specimen temperature of 400 °C but to a much lesser

extent.

B. Metal Wastage Analysis

Weight loss determinations and tube surface examinations by SEM and EDX

techniques were also made on Type 316 and Type 304 stainless steels.

Three specimens, one of Type 316 stainless steel, specimen 4-6, and

two of Type 304 stainless steel, specimens 13-4 and 14-4, were weighed

prior to exposure at tube wall temperature of 590 °C, 590 2 C and 400 3 C,

27
respectively, under the general conditions as described in the first

paragraph of this report. In these cases, after extraction from the

test rig, the entire specimen was rinsed in warm water and subsequently

cleaned in 10% HN0 3 at 50 °C for 20 minute cycles in an ultrasonic bath

until constant weight was obtained. In each case the second cleaning

cycle resulted in no further detectable weight change. The differences

in weight (weight loss) are attributed to loss of metal (wastage) from

the specimen; see Table 1. Fig. 12, 13 and 14 show the stainless

steel tubes after exposure and cleaning. After weighing, sections were

taken from each tube for SEM/EDX analysis.

Fig. 15, a micrograph of specimen 4-6, shows areas of pitting and wastage.

Fig. 16, along with the EDX analysis, indicates regions of cation migration

and concentration. There seems to be no general trend towards con- centration of any metal cation; however, there seems to be a general

trend towards lower nickel concentrations. To see if these conclusions

were reasonable, another area, shown in Fig. 17, was examined. Fig. 18

shows again areas of low nickel. Fig. 18a and 18b, but also regions of

high nickel, Fig. 18c, along with areas of other metal cation concentration. However, the high nickel regions are generally isolated and

not in sufficient number to account for all the low nickel regions.

Although the wastage (weight loss) of this sample was minimal, the nickel

concentration tends to be lower than that of the bulk stainless. It is

possible that this occurs because the nickel is being attacked and lost

through some volatile chemical reaction or it may be that there is a formation of a higher chromium, iron layer that has migrated upwards

covering the nickel. In either case, lower nickel would be observed.

In samples 13-4 and 14-4 a comparison evaluation can be made of the same type of stainless steel exposed at two different temperatures. The

greatest change in weight (wastage) was observed in sample 13-4, while

a minimal amount was lost in sample 14-4; see Table 1. Fig. 19 exhibits

an area of sample 13-4 riddled with holes and depressions between grains.

Fig. 20, 20a - 20f shows metal cation concentration or loss in this area.

In sample 14-4 grain boundaries and grain forms are evident; see Fig. 21.

However, there are not as many pores or holes showing. Black spot regions,

though, are prevalent on the grain surfaces. Fig. 22 shows these areas rather vividly; SEM/EDX analysis indicates them to be areas of low chromium,

low nickel and high iron concentration associated with potassium and sulfur.

These areas are probably the beginning of reaction sites on the stainless

steel surface. Fig. 23 is a high magnification micrograph of one such

spot. EDX analysis, Fig. 23a and 23b, indicates low chromium and low

nickel along with potassium and sulfur in that spot with high metal cation

concentration surrounding blackish area.

A random sample of 304 stainless steel was selected for SEM/EDX analysis

and used as a comparison to the specimens placed in the corrosive

environment. Prior to sectioning, the sample was cleaned with 10% HN0 3

28
at 50°C for 20 minute cycles in an ultrasonic bath. A comparison

of this surface to the surfaces exposed to the corrosive environment

shows that, while grain boundaries are evident in both cases, figures

24-26, intergranular and surface attack seems to be occurring in the

latter cases. EDX examination of the surface of the unexposed specimen

revealed no non-uniform distribution of nickel, chromium or iron.

PI ans : Specimens of Type 316 stainless steel will be evaluated under

oxygen rich conditions with a mixture of 80% by weight K2 C0 3 and 20%

by weight K2 S04 . These specimens will be analyzed by SEM/EDX techniques.

Further gravimetric and SEM/EDX analysis will be conducted on Type 304

stainless steels.

29
Specimen

Specimen

Seed

Gas

Stream

Specimen

Weight

Weight

Change

No.

Stainless

Steel

State

Temperature

In

Grams

In

Grams

In

Weight

Type

°C

Before

Test

After

Test

In

Grams*

and

Cleaning oi

COocr> LO

lo

co

03I

03

CO

p-. o

CD

03Ocr> LO

^3" oCOI CO

COoo• oI oLO

co. LO

03

px Px CO

CNJ LO LO

^3" CO 00 • . • CO CT1 LO

P». 03 03oo

*D4— • -C JZ JZ O E a CJ o CO

•p— •r— JZ CD cc a: +-> S- cd +j CO CO CO c oi o o O CD i

OI

co i

CD

CD

+-> X!

*3" «xP <3-

O CD o o O -M T3

oo GO CO CD CO CJ CJ CD CD JZ oi

+-> CD

^3" OCOI «D"

CDcoa

cd01

LO

px

CO oLO

CO

-ac

Z5o

ico

a)0

QJC.

01

CD

4- -C O +->

+->

O

co +-> jz

+-> -a

CD -a +-> CD u

+> a> O *r-J C JO3

CD C/1 -O

OI -a •<-

zs •> 0 CD jz i

01 -a

-a

4-> *r-E *

30
Fig. 1. Section of Type 316 stainless steel tubing after exposure to

K2SO4 seeded oxygen rich hot gas stream. Note formation of

deposit on the upper surface. Tube temperature 590 °C.

Fig. 2. Section of Type 316 stainless steel tubing after exposure to

K

2 SO 4 seeded oxygen rich hot gas stream. Note formation of

deposit on the upper surface. Tube temperature 500 °C.

31
Fig. 3. Section of Type 316 stainless steel tubing after exposure to

K 2 SO

4

seeded oxygen rich hot gas stream. Note formation of

coating on the upper surface. Tube temperature 400 °C.

Fig. 4. SEM micrograph, 20 X, of K

2 SO4 deposit on leading surface of Type 316

stainless steel. Specimen temperature 590 °C^gas stream oxygen rich.

32
Fig. 5 . SEM micrograph, 200 X, of gas stream-depost interface showing

potassium rich band.

Fig. 6 . SEM micrograph, 500 X of Type 316 stainless steel -deposi t interface.

Specimen temperature 590 °C, gas stream, oxygen rich.

33
Fig. 7. SEM micrograph, 2000 X, of center region of Fig. 6 showing reaction

area at interface. Lettered regions correspond to labeled EDX

spectra, Figs. 7a-7h. Specimen temperature 590 °C, gas stream

oxygen rich.

S K Cr Fe Ni

Fig. 7a. EDX spectra of regions A and B of Fig. 7 showing low Ni concentration

areas. Specie positions and designations on all spectra in this section

of this report are as indicated above.

34
Fig. 7b. EDX spectra of regions C and D of Fig. 7 showing low Cr and low

Ni concentrations in the reaction zone.

Fig. 7c. EDX spectra of regions E and F of Fig. 7 showing high Ni and low

Cr concentrations.

35
Fig. 7d. EDX spectra of regions G and H of Fig. 7 showing high Ni at the

metal edge.

Fig. 7e. EDX spectra of regions I and J of Fig. 7 showing Fe penetration into

the deposit.

fe

36
Fig. 7f. EDX spectra of regions K and L of Fig. 7 showing Fe and Cr penetration

into the deposit.

Fig. 7g. EDX spectra of regions M and N of Fig. 7 showing again Fe and Cr

penetration into the potassium-sulfur deposit.

37
Fig. 7h. EDX spectra of regions 0 and P of Fig. 7 showing penetration of all

three species, Fe and Cr and Ni into the deposit.

Fig. 8. SEM micrograph, 2000 X, of Type 316 stainless steel-deposit interface.

Lettered regions correspond to labeled EDX spectra, Figs. 8a-8e.

Specimen temperature 590 °C, gas stream oxygen rich.

38
Fig. 8a. EDX spectra of regions A and B of Fig. 8 showing low Ni and low Cr

concentrationsat the metal edge.

Fig. 8b. EDX spectra of regions C and D of Fig. 8 showing low Ni and low Cr

concentrations at the metal edge.

39
Fig. 8c. EDX spectra of regions E and F of Fig. 8 showing metal cations in

the reaction zone of the deposit.

Fig. 8d. EDX spectra of regions G and H of Fig. 8 showing metal cations in

the reaction zone of the deposit.

40
Fig. 8e. EDX spectra of regions I and J of Fig. 8 showing metal cations in

the reaction zone of the deposit.

Fig. 9. EDX spectra of regions K and L of Fig. 8 showing Fe farther back

in the deposit.

41
Fig. 10. SEM micrograph, 2000 X, of Type 316 stainless steel-fume deposit

interface. Lettered regions correspond to labeled EDX spectra

Figs. lOa-lOc.

Fig. 10a. EDX spectra of regions A and B of Fig. 10 showing distribution of

Cr, Fe and Ni at the stainless steel edge.

42
Fig. 10b. EDX spectra of regions C and D of Fig. 10 showing distribution of

Cr, Fe and Ni at the stainless steel edge.

Fig. 10c. EDX spectra of regions E and F of Fig. 10 showing absence of Ni in

the reaction zone.

43
Fig. 11. SEM micrograph, 2000 X of Type 316 stainless steel -deposi t interface.

Lettered regions correspond to labeled EDX spectra Figs, lla-llh.

Specimen temperature 500 °C gas stream oxygen rich.

Fig. 11a. EDX spectra of regions A and B of Fig. 11 showing slight loss of

Hi at stainless steel edge as compared to the bulk.

44
Fig. lib. EDX spectra of regions C and D of Fig. 11 showing high Fe and

low Cr and Mi concentrations in reaction zone.

Fig. 11c. EDX spectra of regions E and F of Fig. 11 showing the penetration

of Fe in the deposit.

45
Fig. lid. EDX spectra of regions G and H of Fig. 11 showing concentrations of

Fe, Cr and Ni in the reaction zone of the stainless steel -deposit

interface.

Fig. lie. EDX spectra of regions I and J of Fig. 11 showing penetration of Fe,

Cr and Ni in the deposit.

46
Fig. Ilf. EDX spectra of regions K and L of Fig. 11 showing concentration of

Fe, Cr and Ni at the metal edge.

Fig. llg. EDX spectra of regions M and N of Fig. 11 showing low concentration of

Fe in reaction zone.

47
* u •

.W'' [P . s

/.

'* *.<«/•' V, A /

r

: , 1 , 1 1 1 . l 1 1 1 1 1 1 I 1 1 1 1 1 1

:

Fig. llh. EDX spectra of regions 0 and P of Fig. 11 showing concentrations of

Fe, Cr and Ni in the reaction zone.

Fig. 12. Type 316 stainless steel after exposure for 4 hours at 590 °C

to K 2 SO4 seeded, oxygen rich hot gas stream and subsequent cleaning.

Approximately 2 X.

48
Fig. 13. Type 304 stainless steel after exposure for 4 hours at 590 °C to

K2SO4 seeded, oxygen rich hot gas stream and subsequent cleaning.

Approximately 2 X.

Fig. 14. Type 304 stainless steel after exposure for 4 hours at 400 °C to K 2SO4

seeded, oxygen rich hot gas stream and subsequent cleaning. Approximately

2 X.

49
Fig. 15. SEM micrograph, 400 X, of Type 316 stainless steel exposed for 4

hours at 590 °C to K 2 SO4 seeded, oxygen rich hot gas stream.

Fig. 16. SEM micrograph, 1600 X, of mid region of Fig. 15 showing surface

attack.

50
Fig. 17. SEM micrograph, 400 X, of Type 316 stainless steel exposed for

4 hours at 590 °C to K2 SO4 seeded, oxygen rich hot gas stream.

Fig. 18. SEM micrograph, 1600 X, of mid region of Fig. 17 showing surface attack.

Labeled regions correspond to labeled EDX spectra. Figs. 18a-18c.

51
Fig. 18a. EDX spectra of regions A and B of Fig. 18 showing low Ni and high

Cr concentrations.

Fig. 18b. EDX spectra of regions C and D of Fig. 18 showing high Ni and high

Cr concentrations.

5?
Fig. 18c. EDX spectra of regions E and F of Fig. 13 showing low Ni and high

Fe concentrations.

Fig. 19. SEM micrograph, 450 X, of Type 304 stainless steel exposed for 4

hours at 590 °C to a K2 SO4 seeded, oxygen rich hot gas stream.

53
Fig. 20. SEM micrograph, 1900 X, of the mid region of Fig. 19 showing surface

attack. Labeled regions correspond to labeled EDX spectra, Figs.

20a-20f.

Fig. 20a. EDX spectra of regions A and B of Fig. 20 showing Ni depleted regions.

.

54
Fig. 20b. EDX spectra of regions C and D of Fig. 20 showing another Ni

depleted region.

Fig. 20c. EDX spectra of regions E and F of Fig. 20 showing a region of high

Cr concentration.

55
Fig. 20d . EDX spectra of regions G and H of Fig. 20 showing another region

of high Cr concentration.

Fig. 20e. EDX spectra of regions I and J of Fig. 20 showing another region

of high Cr concentration.

56
Fig. 20f. EDX spectra of regions K and L of Fig. 20 showing regions of high

Cr and low Ni concentrations.

Fig. 21. SEM micrograph, 350 X, of Type 304 stainless steel specimen 14-4,

exposed for 4 hours to a K2SO4 seeded oxygen rich hot gas stream.

57
Fig. 22. SEM micrograph, 375 X, of another region specimen 14-4 showing

what appears to be the onset of corrosion.

Fig. 23. SEM micrograph, 1500 X, of the mid region of Fig. 22. Labeled

regions correspond to labeled EDX spectra, Figs. 23a and 23b.

58
Fig. 23a. EDX spectra of regions A and B of Fig. 23 showing nickel depletion in

a dark region.

Fig. 23b. EDX spectra of regions C and D of Fig. 23 showing a Ni depleted region

C and a region of high Ni and high Cr concentration.

59
Fig. 24. SEM micrograph of Type 304 stainless steel, unexposed but cleaned.

Note grains and grain boundaries , 1600 X.

Fig. 25. SEM micrograph, 1900 X, of Type 304 stainless steel after exposure

for 4 hours at 590 °C to a K2SO4 seeded hot gas stream and subsequent

cleaning. Note grain boundaries and grain surface.

60
Fig. 26. SEM micrograph, 1400 X, of Type 304 stainless steel after exposure

for 4 hours at 400 °C to a K2SO4 seeded hot gas stream and subsequent

cleaning. Note grain boundaries and grain surface.

61