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Distribution Category UC-90g

DEVELOPMENT, TESTING AND EVALUATION OF

MHD-MATERIALS

for the period April - June 1977

H. P. R. Frederikse, T. Negas and S. J. Schneider

Inorganic Materials Division

Institute for Materials Research

National Bureau of Standards

Washington, D. C 20234

Date Published: June 30, 1977

PREPARED FOR THE UNITED STATES

ENERGY RESEARCH AND DEVELOPMENT ADMINISTRATION

Under Contract No. E(49-1)-3800

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

States Government. Neither the United States nor the United States ERDA, nor any of their employees, nor any of their contractors, subcontractors,

or their employees, make 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 private owned rights."

Quarterly Report

J

Table of Contents Page

ABSTRACT 1

OBJECTIVES AND SCOPE OF WORK 2

SUMMARY OF ACHIEVEMENTS 3

Task G. Program Management and Coordination 5

1 . 2. Program Review and Consultation Activities

US/USSR Cooperative Program

Task I

.

Operational Design Properties 7

1 . 2.

3.

Viscosity of Coal Slags

Electrical Conductivity

Vaporization and Dissociation Studies on K^SO^

Task J. Corrosion and Diffusion 22

1 . Seed-Slag Interactions

a. Compositional Variations in KAlC^-SiO-, Based Materials

b. Thermochemistry of the CaO-K^O-Al^O^-SiO^ System

c. Thermochemistry of the MgO-K^O-A^O^-SiC^ System

d. Thermochemistry of the FeO -K^O-A^O^-SiC^ System

e. Implications of Research on Seed-Slag Interactions

2. Diffusion in Insulator-Electrode Systems

Task K. Materials Testing and Characterization 51

1 . 2.

3

.

4.

5.

Introduction

Structural Analysis of powdered and solid ceramics

Preparation for Future Tests (U-02 Phase III, AVCO)

Testing and Post-Analyses of Electrodes tested at MIT

Examination of Preheater materials and Fly-ash

Task L. Assessment of Steam Plant Components 69
}
1

ABSTRACT

Electrode design and fabrication as well as post-test characterization

received much attention during this reporting period. Several new methods

for joining the ceramic to the metal base are being considered, including a thick cermet layer and a mechanical attachment. A series of 11 anodecathode pairs (5 different designs, 3 different materials) are being fabricated (with outside assistance) for a test in the AVCO generator under

slagging conditions.

An appreciable amount of time has been spent in consultation with

Westinghouse , Battelle and APS Materials Inc., in connection with the

planning of Phase III of the U-02 testing program.

A number of electrode/insulator modules that have been tested in the

MIT rig were analyzed at NBS (ANL, GE and NBS designs)

The large volume of viscosity data, collected over the last 4 years,

are presently being analyzed and presented in the form of a comparison with

a "base slag composition" (half-way between Illinois #6 and Rosebud)

Previous data of Halstead on the vapor pressure of K2 SO4 together with

results of new NBS measurements have been carefully compared and evaluated.

The saturation pressures of K^SO^ appear to be about 5 times lower than

previously assumed, which has important implications for seed-slag separation

procedures

In the assessment of alloys for downstream components attention has been

focussed on mechanical strength and resistance to H^S-attack.

2

Objectives and Scope of Work

The overall objective of this program is to obtain chemical and physical

definition of high temperature materials which have shown promise for use

in coal-fired open-cycle MHD power systems. Major problem areas in which

investigations will be concentrated are:

1. Characterization of coal slag and its effects on system components

and performance at prototype temperatures.

2. Development of electrode materials which provide adequate performance

over extended periods of time.

3. Insulating materials which limit thermal losses and are resistant

to prolonged thermal and erosion effects.

4. Preheater materials which can withstand the operating modes of

separately and directly fired operation.

5. Seed recovery methods from slag which are technically and economically

feasible

.

6. Phase equilibria and diffusion rates of seed in slag and the corrosive

action of combination on system components and materials.

7. Durability of prototype MHD sub-systems.

The program is designed to contribute to the solution of these problems by

providing much needed data on candidate materials and by evaluating test sam- ples and structures that have been subjected to real or simulated MHD conditions.

The activities are grouped under six tasks:

G. Program Management Coordination (Assisting ERDA in coordination, planning

and review of the various MHD-Materials Development Programs)

.

H. U-02 Materials Testing and Characterization (Coordination of U-02 Test

Activities, Phase I). (Terminated June 30, 1976).

I. Operational Design Properties (viscosity, electrical conductivity,

vaporization)

.

J. Corrosion by Seed and Slag (phase equilibria, diffusion)

.

K. Materials Testing and Characterization (test coordination, pre- and

post-test analysis)

.

L. Assessment of Steam Plant Components (corrosion resistance of metals

and alloys)

.

3

Summary of Achievements (April-June 1977)

(completed Milestones-See Work Statement for FY 1977)

Third Quarter accomplishments in terms of the projected milestones

of the NBS-ERDA contract are listed below. ERDA directives requesting

change of direction or emphasis (electrodes for U-25, characterization

of electrodes for other contractors, etc.) have necessitated some changes

in the chronological order of projected targets.

lc. Planning meetings and consultation sessions have been held to prepare

for the U-02 materials test (Phase III) to be conducted in early 1978.

2. A number of NBS staff members particpated in project reviews, preparation

of program plans, etc.

4a. Measured the viscosity of a New Hampshire flyash (AVCO) and a, number of

synthetic slags.

10b. Determined first versions of the phase diagrams for the quaternary

systems K^O-Al^O^-SiO^-MgO and K^O-Al^O^-SiO^ -CaO.

13a. Measured the K^SO^ vapor pressure over K^SO^

19a. The electrical conductivities of a [60% spinel (MAFF-31 ) -40% Cr] - cermet

and of a (not-so-insulating) sample of SrZrO^ were determined.

30a. Measurements have been initiated on diffusion of iron from Mg-Al-Fe

spinel into MgO.

39b. Electrode/insulator modules were designed, fabricated (outside NBS)

and prepared for testing in various MHD test facilities (MIT, AVCO)

39c. Samples of 35 materials, that are being used by ERDA contractors for

MHD-electrodes , insulators and preheaters, have been characterized at NBS

by means of X-ray diffraction with respect to their phase composition.

40c. Optical and S-E-Microscopy as well as X-ray diffraction have been used

for the evaluation of 3 electrode assemblies (NBS, ANL and GE) tested at MIT.

43a. Compiled a list of mechanical properties of promising alloys in

particular for steam super heater application.

Talks and Papers 4

Coal Slag Properties Related to MHD, Conference on High Temperature Sciences

Related to Open-Cycle , Coal-Fired MHD Systems, Argonne Natl. Lab., 4/6/77,

W. Capps .

Some Vapor Pressure Data on K2 0 Containing Substances, Conference on High

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

National Lab., 4/6/77, E. R. Plante .

Status of MHD Materials, Conf. on High Temperature Science Related to Open

Cycle, Coal-Fired MHD Systems, Argonne Natl. Lab., 4/6/77, S. J. Schneider .

The Second Joint Test of a U.S. Electrode System in the USSR U-02 Facility,

G. Rudins, S. J. Schneider , T. Negas , B. R. Rossing, J. L. Bates,

G. P. Telegin, T. Bordina, 0. Zamislov and V. Zalkind, 16th MHD Symp .

,

Pittsburgh, PA May 16-18, 1977.

Electrodes and Insulators: Design and Materials Considerations, H . P . R . Frederikse and W. R. Hosier, 16th MHD Symp., Pittsburgh, PA, May 16-18, 1977.

Some Properties of Coal Slags of Importance to MHD, 16th MHD Symp., Pittsburgh,

PA, W, Capps , May 16-18, 1977.

Interaction of MHD Seed (K2 0) with Multicomponent Oxide Solids and Liquids,

L. P. Cook , E. R. Plante, C. L. McDaniel, T. Negas, and R. S. Roth, American

Ceramic Society Meeting, Chicago, II., April 25, 1977.

5

Task G. PROGRAM MANAGEMENT AND COORDINATION (S.J. Schneider) 1.

Program Review and Consultation Activities

S. J. Schneider participated in the regular ERDA-MHD staff meetings

and in assorted program review/coordination meetings and briefings. The

meetings included

1. ANL MHD Conference (invited speaker) April 4-6

2. Program review at Montana Tech and Montana State University -May 2-6

3. 16th MHD Symposium (presentation of paper on the US-USSR

Phase II Test - Pittsburgh, May 16-18)

4. Materials Ad Hoc Working Group Meeting (U-02 Phase III) April 7 5. Steering Committee - US-USSR Program

6. Technology Task Force - US-USSR Cooperative Program

The conclusions and recommendations resulting from these meetings are reflected in reports to ERDA, through direct consultation with ERDA staff or through documents published elsewhere.

NBS receives on a regular basis, MHD technical reports and proposals

for review. During this period NBS provided 3 such evaluations.

Part of the NBS experimental program involves characterization of

materials component elements tested in various MHD test rigs and facilities

throughout the United States (and USSR) . In order to expedite the dessemination of this information, NBS has instituted a separate reporting system in

addition to the regular Quarterly submissions. The system consistes of the

following elements: 1. Telephone report to submittor after initial

examination, 2. Follow-up letter report, and 3. Detailed report (if warranted)

to submittor and ERDA. The detailed report will be distributed to general

contractors on an as needed basis as determined by ERDA; however highlights

of the detailed report will be abstracted and published in the NBS Quarterly

reports. In addition, general overall summary reports will be compiled from

time to time and published in appropriate journals and special editions.

Inquiries regarding the status of any particular characterization or requests for additional assistance should be directed to:

A. Perloff 301-921-2900

alternates

T. Negas 301-921-2843

W. Hosier 301-921-2940

2 . US-USSR Cooperative Program

a. U-25 Electrode Program

Because of its importance this activity will have, in the coming

months, renewed emphasis. NBS responsibilities and accomplishments are reported elsewhere in this report.
.
6

b. US-USSR Materials Program Activities

During this reporting period the US-USSR Ad Hoc Materials Working Group

held several informal meetings concerning the implementation of the Phase III

test at the U-02 in Moscow. The overall activity is being managed by

Westinghouse; other participating groups include NBS , Battelle NW, GE , and

ANL. The present schedule calls for proof-testing of candidate electrode

systems at the Westinghouse facility in August-September 1977, delivery of

the finished electrode wall to Moscow in November-December 1977 and actual

testing in January 1978 (note scheduled dates are tentative and may be

changed)

.

Upon the completion of the Phase III test the Materials Working Group

will cease its activities as a separate entity; its functions will be

incorporated into the efforts of the newly formed Task Force on MHD

Technology Development (US-USSR)

.

In addition to the above, the Joint US-USSR chapter on materials

(Status Report) has been again revised to incorporate new information

provided by the USSR. It is believed that this will be the last version

and that the chapter is "ready to go to press."

7

Task I. OPERATIONAL DESIGN PROPERTIES

1.

Viscosity of Coal Slags (W. Capps and D. A. Kauffman, Inorganic Glass Section)

Introduction

Part of the program of characterizing MHD materials is the preparation

of synthetic slags and the determination of viscosity of these and of any

real slags furnished to NBS by other contractors.

Slag Preparation and Viscosity Measurement

1. A power plant flyash from a New Hampshire coal was sent to NBS by

Mr. Stan Petty, AVCO-Everett Research Laboratory. A melt was made and

its viscosity was determined and is shown in Fig. 1 along with the analysis

furnished by AVCO.

2. Twenty synthetic slags from the NBS studies were sent to Dr. James Anderson

at Montana State University for use in his program. Three specimens were

sent to Mr. Harry Rausch at the G. E. Valley Forge Space Center.

3. Systematic additions were made to NBS Base Composition, K-884, to

determine the viscosity dependence on concentration of silica, alumina,

iron-oxide, etc. Preliminary results are shown in Figs. 2, 3 and 4.

Findings to date indicate (1) a nearly linear, large, positive dependence

on silica concentration, (2) a small negative dependence on alumina, and

(3) a large, negative, nearly linear dependence on the amount of iron

oxide. Although opposite in sign, the rate of change due to iron is the same as the silica rate.

4. Two talks were given: "Coal Slag Properties Related to MHD" was given

at the High Temperature Chemistry Symposium at Argonne National Laboratory

in April and "Some Properties of Coal Slags of Importance to MHD" was given

at the 16th Symposium on Engineering Aspects of MHD at the University of

Pittsburgh in May. An abstract of the talk will be printed in the proceedings .

Future Work

The series described in paragraph 3 above will be continued to determine

the effects on viscosity of calcium and magnesium oxides.

This new data together with an extensive bank of earlier data will be

used to improve equations and/or develop a suitable graphical method for

purposes of prediction of viscosity of coal slags from ash analyses.

*NOTE : Compositions of melts referred to in Figures 2, 3 and 4 appear

in Table I.

8

Table 1

Coal Slag Compositions (wt. %)

NBS

Melt No. sio

2 Al 2°3 Fe 2°3 CaO MgO

K-884 47.77 21.76 16.58 10.57 3.32 Base

K-887 52.77 19.68 14.99 9.56 3.00 +5% SiC>

2 K-901 57.77 17.59 13.41 8.55 2.68 +10% "

K-913 42.77 23.84 18.16 11.58 3.64 -5% "

K-915 37.77 25.93 19.75 12.60 3.96 -10% "

K-921 45.33 25.76 15.73 10.03 3.15 +5% Al 0. K-926 50.83 16.76 17.63 11.25 3.53 -5% "

K-931 53.88 11.76 18.70 11.92 3.74 -10% "

K-933 44.90 20.45 21.58 9.94 3.12 +5% Fe

2 0

K-934 42.04 19.15 26.58 9.30 2.92 +10% "

K-953 39.18 17.84 31.58 8.67 2.72 +15% "

K-955 50.63 23.06 11.58 11.20 3.52 -5%

K-958 53.49 24.36 6.58 11.84 3.72 -10% "

9

1200

1300

1400

1500

1600

1700

TEMPERATURE

(C)

10

O

UJ

or

Z>

<

(T

UJ

a.

2

UJ

l—

<

o

-J

cn

u_

o

O

UJ

Ll

UUJ

OJ

cfi

U_

11

to

01

OJ

O

O

NO

O

CD

O

O

in o

LU

o:3

V~

o <

o cr

lu — CL

UJ

H

O

O

to

o

o

CVJ

(sasiod B 0 |) A1IS00SI

A

Fig.

3.

EFFECT

OF

ALUMINA

12

o

o

ro

o

CD

O _ 0 /j

i£2 ii

LJ

tr

Z>

H

<

CK

LsJ 1 K.

LJ

hO

C

ro

O

o

o

o

CM

CP

iZ

EFFECT

OF

IRON

OXIDE

13

2. Electrical Conductivity (W. R. Hosier and A. J. Armstrong)

LaCrOo. After the electrical conductivity measurements on arc plasma sprayed

LaCr03 + 5% MgO were completed (see Jan. -March 1977 quarterly report)

small hair-like crystals were seen to have grown between the electrical

contacts. These crystals were removed and examined by scanning electron

microscopy and energy dispersive x-ray (EDX) analysis.

The crystals were of two types: small rods (0. 5-5.0 ym diameter) and

larger crystals (10-20 ym diameter) which were acicular in appearance.

Both types of crystals were approximately the same composition as determined

by EDX analysis. They were mostly composed of Mg and Cr with a few percent

of Si and Ni present. The presence of silica is detrimental to the

electrical conductivity and high temperature strength of LaCrO-j. The source of these impurities should be eliminated in order to provide the

best possible material for use as an MHD electrode.

Cermet . A sample of MAFF-31 ( 3 MgAl204 : FejC^) + 40% Cr cermet was sprayed in a free stand by APS Materials Inc., Dayton, OH. A portion of this material was tested at several temperatures between 700 °C and 1500 °C. After each high

temperature treatment which lasted for from two to six hours, the sample was cooled to room temperature. Electrical conductivity measurements were made

at the high temperature and then at room temperature to determine the

deleterious effect of oxidation. Conductivity measurements were made using

very high current densities (15A/cm2)

.

Although measurable oxidation is known to occur in Cr metal at

temperatures above 900 °C in air, no significant change was seen in room

temperature electrical conductivity until the sample had been subjected

to a temperature of 1215 °C. Room temperature values ranged from 1000 to

600 ohm- lcm"l after treatment at temperatures up to 1133 °C. After treatment at temperatures between 1215 °C and 1535 °C, room temperature electrical

conductivity values fell from 126 to 55 ohrrrlcm"l. (Fig. 5) .

This Cr-MAFF-31 cermet has been used in several electrode designs as

a lead-out material for MAFF-31. The chromium provides the necessary

electrical conductivity at temperatures below 600 °C as well as a lower

thermal conductivity than the Hastelloy B metal previously used.

Electrical conductivity values at elevated temperatures ranged between

300-400 ohm-

-*-cm- l at temperatures between 730 °C and 1468 °C. At 1535 °C

a sharp decrease in electrical conductivity was seen (35 ohm- lcm-

^)

.

After removal, the sample still possessed considerable strength and

there were no apparent cracks despite repeated thermal cycling.

SrZrO ,, . A sample of SrZrO^ designated SZ-CE0104 was received from Westinghouse

Research and Development Lab., Pittsburgh, PA. This was a sintered piece

which was single phase material of good crystallinity but appeared slightly

pink in color.

14

The electrical conductivity of the sample was measured as a function of

temperature and several partial pressures of oxygen. The results are shown

in Fig. 6. The electrical conductivity of this sample would be too high to

be effective as an insulator material for high temperature use in an MHD

generator. Although the data is in general agreement with that given in

" Introduction to Ceramics V by W. D. Kingery, H. K. Bowen and D. R. Uhlmann,

there is the possibility of the high conductivity being due to impurities

present. So far no chemical analysis was made.

Another sample of known high purity has been received and measurements

are in progress.

i

\JI_yOV_/

f

I

VTT'T

^OTTTT-unT;

15

T(K)

Fig. b . Crosses represent the conductivity of the cerment at high temperatures.

Circles indicate the room temperature conductivity after treatment

at the temperature shown on the abscissa.

CONDU

C!

IV

IT

Y(ohm-cm)

16

T(K)

Fig. 6. Electrical conductivity of a sample of SrZrO^ from Westinghouse

(SZ-CE0104) . Sintering time, temp. & atmos. not given.

17

3. Vaporization and Dissociation Studies on K SO (E. R. Plante)

Z. 4t The vaporization program has been involved in the determination of

K(g) pressures over various K^O-containing materials. The purpose of

these measurements has been to provide data which can be used for

predicting the equilibrium concentration of K^O in slag solutions and

to predict the temperatures and K(g) pressures at which MHD structural

materials will resist corrosion.

Although the molecule ^SO^g) does not enter directly into these

considerations, it is one of the more important K bearing molecules at

temperatures from 1600 - 2200K and plays a significant role in defining

the activity of K^O in the vapor phase as well as the condensation

temperature of K2 SO4

. A comprehensive review article on the vaporization

thermodynamics of K^SO^ by Eliezer and Howald (1) was recently presented.

Their critical evaluation of the existing data concluded that the most

reliable measurements were probably those of Halstead (2)

;

a viewpoint

with which we and most of the MHD community probably agree. However,

the lack of reliable measurements confirming his data and a review (3)

of the vapor pressures of S^O-containing substances important to MHD

made it appear that some measurements to* confirm his results were in

order

.

Halstead's measurements were made using the gravimetric Knudsen

method below, and the transpiration method above the melting point. (1349 K)

His Knudsen data were obtained using cells pressed from equimolar

mixtures of Ai^C^ and MgO. He recognized potential problems due to

cell porosity and/or reaction of the with the cell. Measurements

were therefore made using cells in which the inner or outer surface

was covered with Pt. Since these latter results agreed with the former,

it was concluded that the Knudsen cells were both non-porous and unreactive

to ^SO^c). The transpiration results were not regarded by Halstead

to be as trustworthy as the Knudsen results because of difficulties in

obtaining vapor pressures independent of flow rate at higher temperatures

(1668K) . However, within the experimental error, the vapor pressure

data were in reasonable agreement with a vapor pressure equation based

on the Knudsen data extrapolated into the liquid range by correcting for

the change in slope and intercept due to the heat and entropy of fusion.

While Halstead recognized the potential errors due to porosity of

the Knudsen cell and/or reaction with Al^O^ or MgO, the presence of

Al2®2 is a common factor in both his transpiration and Knudsen experiments.

Although it is not known whether the rate of reaction would be sufficiently

high to cause serious error it can be predicted from data in the JANAF (4)

tabulation and some measurements previously reported from this laboratory (5)

that the reaction of K2 SO4 with A£ 2°3 to form ^0 • 9A& 2O 3 is thermodynamically

favorable. These are additional reasons why new measurements of the

vapor pressure of K2 SO4 were considered worthwhile.

There is no doubt that the overall vaporization behavior of K2 SO4 can be described in terms of the sublimation or vaporization reaction,

18

K

2 S04

(c, ft,) = K

2 S0

4

(g)

and the dissociation reaction,

K

2 S04 (c, Z) = 2K(g) + S02 (g) + 0

2 (g)

The existence of K S0

4

(g) molecules has been demonstrated through

previous mass spectrometric measurements (6, 7) as well as some recent

observations in our own laboratory and indirectly by the observation

that the measured pressure is greater than that predicted for the

decomposition pressure alone. It is worth noting that decomposition

pressures involving well characterized vapor species and condensed

phases can generally be calculated from thermodynamic data more

accurately than they can be measured. The major question therefore

concerns the relative importance of the vaporization and dissociation

reactions. Concentrations of other species in the gas phase such

as K

2 0, KO, and SO^ can be shown to be minor (< 1%) in comparison.

Experimental Measurements (In collaboration with C. Olson and T. Negas)

Mass effusion measurements were made using the Mettler Microbalance

System. The K

2 S0

4 sample was NBS-ICTA Standard Reference Material 760.

The effusion cell was constructed of Pt and had an effusion orifice

.020" in diameter. Weight loss data were obtained at fixed temperature

and recorded in the form weight vs time. Linear portions corresponding

to m/t were selected and used to calculate apparent pressures. The

rate of vaporization observed corresponds to that due to effusion of

K(g) , 0

2 (g) and SC> 2 (g) as well as K

2 S04

(g) . The observed rate of

evaporation can bemused to calculate an apparent K

2 S04

(g) pressure using

the equation

211 RT

M

in which m/t is the weight loss per unit time, a is the area of the

effusion orifice, c is the Clausing factor, R is the gas constant,

T is the absolute temperature and M is the molecular weight of K^SO^.

Comparison of the apparent pressure with the calculated dissociation

pressure shows that a correction for effusion of dissociation products

is necessary.

The contribution of dissociation products to the apparent K

2 S0

4

(g)

pressure can be shown to be

5 pk

2 so4 - V2 V

M

K2 so/‘\

Here, M^ and M

K

are the molecular weights, and P is the potassium

pressure^frdm the dissociation reaction under Knudsen flow conditions

which can be determined from the relationship,

1/4

PK

= ( Keq/- 290 >

19

in which Ke^is the equilibrium constant for the dissociation reaction.

Values of P K were obtained at 100K intervals using data from the JANAF (4)

tables. Least square lines for use above or below the melting point

were obtained using 3 calculated pressures at 100K intervals above or below the melting point respectively. The data are given in table 2 which lists the temperature in K, the apparent K^SO^ pressure in atmospheres,

the corrected K

2 SO4 pressure in atmospheres, and the third law heat of

sublimation at 298K. The third law he<_ts were calculated using the

recent Gibbs function data made available by Gurvich et al ( 8 )

.

The second law heat of sublimation calculated by least squares using

the method of Cubicciotti (9), the above Gibbs functions, and based on

data below the melting point is 81140 ± 1960 cal mol

- '1

' compared to the

average third law heat of 81780 ± 60 cal mol“~ indicating good agreement

between the entropies of K2 S 0

4

(g) (45.313 estimated from molecular

constant data by Gurvich ( 8 ) vs 45.820 ± 1.547 determined by the second

law method at 298K) . The errors quoted are standard errors. The standard

heat of formation of K-SO. (g) at 298K is - 261760 cal mol - -1- based on our

2 4 -1 average third law heat and the JANAF (4) value of -343540 cal mol for

the standard heat of formation of K^SO^ (c) . ERDA contractors wishing to

use the standard heat of formation of K-^SO^tg) given here should also use

the Cp data for K^SO^Cg) given by ( 8 ).

Careful analysis of the data points above the melting point show

that they are incorrect because of unknown systematic error (they give

rise to a negative heat of fusion for example) so they were excluded

from consideration in calculating the second and third law heat.

According to our data and analysis the saturation pressure of

K2 SO^ is about 1/5 the value measured by Halstead. This may have

important implications in design considerations for seed-slag separation

in MHD systems. As a rough guess, we estimate that this will increase

the condensation temperature of K

2 SO4 in an MHD system by about 50K.

Future Work:

During the next quarter measurements on the K pressure over K2 0-Si02 _ A 5- 2 O

3 solutions will be continued. Additional measurements on K^SO^ using

a larger effusion hole are planned in order to extend the temperature range • in which measurements of the vapor pressure below the melting point are possible and to confirm the present results. Measurements by mass spectrometry on KFeO^ will be attempted.

20

TABLE 2

K^SO^ Vapor Pressure Data

T°K P apparent

1271 5.01-6

1316 1.25-5

1271 4.00-6

1196 7.75-7

1221 1.44-6

1246 2.38-6

1271 4.40-6

1296 7.27-6

1321 1.28-5

1196 7.27-7

1221 1.29-6

1246 2.18-6

1296 7.27-6

1321 1.20-5

1345a 1.94-5

137 0a 3.18-5

1351 a 2.23-5

1319 1.06-5

1270 3.90-6

1396a 5.15-5

1420a 8.24-5

(atm) P corrected

3.98-6

9.65-6

2.97-6

6.16-7

1.14-6

1.81-6

3.37-6

5.44-6

9.62-

6

5.68-7

9.87-7

1.61-6

5.44-6

8.82-6

1.41-5

2.32-5

1.63-

5 7.55-6

2.89-6

3.74-5

6.05-5

(atm) AH°

29 g (cal)

81267

81383

82007

81572

81566

81851

81688

81821

81650

81765

81906

82141

81821

81878

81847

81617

81729

82182

82020

81423

81091

a Above melting point, not used in second or third law analysis.

21

References

1. I. Eliezer and R. A. Howald, J. Chem. Phys . , 65_, 3053 (1976).

2. W. D. Halstead, Trans. Farad. Soc . , 66_, 1966 (1970).

3. E. R. Plante, Proceedings of the Conference on High Temperature

Science Related to Open-Cycle, Coal-Fired MHD Systems, ANL-77-21

(1977)

.

4. JANAF Thermochemical Tables, Dow Chemical, Midland, Michigan.

5. E. R. Plante, C. D. Olson, and T. Negas, Sixth International Conference

on Magnetohydrodynamic Power Generation, Washington, DC, 211 (1975)

.

6. P. J. Ficalora, 0. M. Uy, D. W. Muenow, and J. L. Margrave, J. Amer

.

Ceram. Soc., 51 , 574 (1968).

7. L. N. Gorokhov, High Temperature Technology, Proceedings of the

Third International Symposium on High Temperature Technology,

Butterworths , London, U. K. (1967).

8. L. V. Gurvich, O. V. Dorofeeva and V. S. Yungman, Proceedings of the

Conference on High Temperature Science Related to Open-Cycle, Coal-Fired

MHD Systems, ANL-77-21 (1977).

9. D. Cubicciotti, J. Phys. Chem., 70, 2410 (1966).

22

TASK J. CORROSION AND DIFFUSION PROCESSES

1 . Seed-Slag Interactions

a. Compositional Variations in KAlO^-SiO^ Based Materials (L. Cook, W. Brower)

To investigate the behavior of KA10 -Si00 based compounds at high

temperatures , container metals suitable for preparation of sealed

capsules in experiments above 1800 °C were sought. Experiments with Ta

and Mo were performed, using (KAIO^) ^ (SiO ) , prereacted at 1600°C.

Capsules were loaded and welded in an inert atmosphere. Experiments were

performed under a forming gas atmosphere in an inductively heated graphite

susceptor with zirconia packing. Temperature reached ^22001100 °C for about

10 min. The tantalum capsule ruptured and showed extensive signs of

reaction (Figure 1) . Reaction products appear to be free of K 0,

although this may have been completely volatilized. The tantalum reaction

product is rich in A^O^ ( aPParently no other elements besides Ta, and

possibly 0) . The reaction has occurred both on the inside and the outside

of the capsule - a thin layer of unreacted Ta is still present between

reaction crusts . Whether the capsule rupture preceded reaction or whether rupture occurred because of reaction on the inner surface is

uncertain. The latter seems unlikely, because there is no strong sign

of reaction adjacent to the melted (KAIO^) ^ (SiO^)

•

Rupture may have

been caused by hydrogen embrittlement rendering che Ta incapable of

withstanding the increased internal pressure at 2200° G - reaction was then possibly caused by contact of (KAIO^) (SiO^) with forming gas.

Further experiments using different atmospheres will test the feasibility

of using Ta.

The molybdenum tubing worked much better in the same experiment

(Figure 2-a)

.

The capsule held, and no significant evidence of reaction

between capsule walls and (KAIO^) _q (S iC^ ) was detected (Figure 2-b)

.

Energy dispersive x-ray spectra of melt snowed no detectable Mo; spectra

of tubing adjacent to melt showed no detectable K. A thin unidentified

skin a fraction of a micron thick was seen adhering to the Mo wall adjacent

to melt in some areas (Figure 2-c)

.

Presumably this is an MoSiO -related

substance. It appears that Mo will be acceptable for sealed high

temperature studies of KAlO^-based materials. Reaction with other com- ponents such as CaO will also be tested.

b. Thermochemistry of the CaO-K^O-Al^O^-SiO^ System (L . Cook)

Experimental work conducted in the last quarter together with

published data have been used to construct the working model of subsolidus compatibilities in the system CaO-K^O-Al^O^-SiO^ , shown in

Figure 3. This system is especially important in understanding the

23

Fig. 1. Fractured surface of Ta capsule with (KA1C> 2

) 79 (SiC> 2 ) 3 Q

sealed inside, run at ^2200 °C in forming gas. Remainder

of melt is visible in bottom (40X)

.

Fig. 2 a. Mo capsule, experiment identical to Fig. 1 . Melt in bottom of capsule (35X)

.

Fig. 2 b. Contact between melt and Mo (1400X)

.

Fig. 2 c. Thin skin adhering to Mo wall (left) at interface

with melt (right) 7000X)

.

25

process of seed capture by western coals , some of which are markedly

more CaO-rich than bituminous midwestern and eastern coals. This latter

point is clear from comparison of Figures 4 and 5, CaO-SiO^-Al^Og

"projections" of multicomponent analyses of bituminous and sub-bituminous

coals, respectively. Also shown on these diagrams are projected tie

lines between phases , taken from Figure 3 . The chemical systems most

important for study, in relation to the compositional distribution of

natural coal ashes from the U.S. are thus apparent. Noteworthy is the

difference in distribution of analyses between the two coal types. CaO

in bituminous coals has a well defined distribution about a maximum

of 't 3 mole % (Figure 4) . CaO in sub-bituminous coals shows no tendency

for a well-defined distribution, with values spread between 1 and 40 mole %

(Figure 5) . This conclusion is reinforced by the "average" analyses of

five Montana coal formations (with standard deviations) , included for

purposes of comparison in Figure 5. In spite of the above difference,

a marked similarity between the two types of coals is apparent in mole %

SiO /(SiO +A1 0 ) ratios (Figure 6). The great majority of analyses fall

between 0.67 (tne projected composition of KAlSiO^) and 0.80 (the projected

composition of KAlSi^O^) , with well defined medians at 't 0.77 and ^ 0.75

for bituminous and sub-bituminous coals, respectively. Referring to

Figure 5, many of the western type coals fall within the projected area

of the system CaSiO -CaAl^Si^Og-KAlSi^Og. It is for this reason that

a study of this system is proposed as an important step in understanding

the behavior of western coal slags in the channel (revised milestone

17-e) . For the bituminous coals, iron is a much more important component

than CaO, as discussed below.

A thermodynamic analysis of the CaO-K^O-Al^Og-SiO^ system is being

conducted in conjunction with experimental determination of invariant

vaporization equilibria in the system and measurement of potassium

activities in melts of invariant composition. This analysis is being

done with the following assumptions

:

a) subsolidus equilibria are as in Figure 3

b) solid solutions involving substitution of K^O do not occur

with the exception of Ca Al^SiO^-KCaAlS^O^, which is assumed

ideal up to the maximum of 20 mole %

c) other solid solutions are not significant in this analysis

d) within each four phase volume, potassia chemical potentials

in the system and potassia "activities" in the gas phase

in equilibrium with condensed phases (defined as PK

2 ’ PO n 1/2

> are fixed by a uniquely defined univariant vaporization2 equation (Table 1) . This is true up to the temperature of

invariant melting and is a reasonable extrapolation for melts

of invariant composition up to plasma temperatures

.

The value of this type of analysis lies in the correlation between the

subsolidus (for which reliable thermochemical data are available) and the

realm of plasma/slag interaction.

CaO

AI2O3

o;n

Fig. 3. Working model of subsolidus equilibria for the system

CaO-K^O-Al^O^-SiO^ (mole %) . Abbreviations as in Table 1.

27
'
28

8502 /(S5©2 + ^

2®3

^

Fig. 5. CaO-A^O-^-SiC^ projection of analyses for sub-bituminous coals,

lignites and unclassified western coals (Bureau of Mines Bulletin

567) in relation to subsolidus equilibria. Data points with bars

show "average" and standard deviations of five Montana coals (after

Diebold)

29

Fig. 6. Mole % silica to silica-plus-alumina ratio in recalculated analyses,

U.S. coal ashes (Bureau of Mines Bull. 567). Grouped according to

bituminous (a) and sub-bituminous or lower (b) rank. Kls and Lc

indicate projected compositions of KAlSiO^ and KAlSi^O^, respectively.

30

Calculated equilibrium curves for invariant melts in three of the

volumes shown in Figure 3, corresponding to reactions 3-a, 4-a, and 7-a

of Table 1 are given in Figure 7. The compositions and invariant temperatures of these melts can only be estimated at present; it is expected

that these will be established during the next quarter. The upper reaction

curve in Figure 7, corresponding to reaction 7-a in Table 1, is nearly

coincident with that calculated for the eutectic melt between ^si 0 , KAlSiO^ and KAlSi^O^, which contains ^ 45 wt % K^O. However a problem

has been encountered in the thermochemical analysis, in that topological

relationships of Figure 3 imply the following order of decreasing yK Q

in the compositional volumes of Table 1: 4, 3, 2, 1, (5, 7), (6,

2 8), 9, 11, 10. Given that

y = y°

K

2 0

yK

2 0

+ RT In Pk

2 o'

a similar relationship should hold among calculated potassium pressures.

The relation among curves in Figure 7 is in the order of decreasing P^ : 7-a, 3,4. It is hoped that measurements during the next quarter, 2 with continued thermodynamic analysis, will resolve this problem.

Reference: F. Diebold, 1977, Characterization of Coal for Open

Cycle MHD Powder Generation System, Project Review,

MERDI, Contract EF-77-C-01-2524

c. Thermochemistry of the Mg0-K

2 0-Al

2 0

2 ~Si0

2 System (L. Cook)

Figure 8 shows a working model for subsolidus compatibilities in

the system Mg0-K

2 0-Al

2 0--Si0 , derived principially from the work of

Schairer (1954) and Lutn (1957) , with limited experimental verification

in this laboratory. The importance of MgO as a component of western- type

coal slags, when recalculated on a mole % basis, is apparent from Figure 9, a Mg0-Si0

2

~Al

2 0_ "projection". MgO content ranges up to 20-25 mole %,

with a maximum at ^ 7 mole %. By contrast, bituminous coals (Figure 10)

usually contain less than 5 mole % MgO, with a maximum at ^1.5 mole %.

The system Mg0-K

2 0-Al2

C> 3

-Si0

2

is thus important primarily in studying

the behavior of western coal slags in the channel, but also in corrosion

studies involving seed, slag and MgAl

2 0 or similar insulators. Nearly

all the sub-bituminous coals project within the Mg

2 Si0^-KAlSi0^-KAlSi2 0^

system, only partially studied by Luth (1967)

A preliminary thermodynamic analysis of the Mg0-K

2 0-Al202 “Si02 system

has been initiated as a prelude to further determination of invariant

equilibria in the system and measurement of potassia activities in melts

of invariant composition. This analysis is being completed with assumptions

similar to a, b, c, d in the analysis of the Ca0-K

2 0-Al

2 0

3 -Si0

2 system

above

Calculated potassium pressures as a function of temperature for

selected invariant melts in four-phase volumes in the systems

Ca0-K20-Al203~Si02 and MgO-K^O-A^C^-SiC^ (dashed) in relation

to seeded plasma (1 wt % K^CO^) curve (supplied by E. R. Plante).

Fig. 7 .

(Mg,FeO)

32

8 . Working model of subsolidus equilibria for the A12°3 _Sl0

2 (mole % ) • Abbreviations as in Table

(M,F) Sp = (Mg, Fe)Al

2 0

4

system (Mg , Fe) 0-K 0-

1 , as well as

(M,F) Sa = (Mg, Fe)

4

Al

10

Si

2 O

23

(M,F) Cd = (Mg, Fe) A1 Si 0

2 4 5 18

01 = (Mg,Fe)

2 Si0

4 = (Fo , Fa)

Pxn = (Mg,Fe)Si0

3 = (En , Fs

)

33

MgO

Flg " 9

' ii°"

Ag2°3 ' gnxtes

S “2 Pro3ecti°n °f analyses for sub-bituminous coals and unclassified western coals - see Fig. 5.



35

One curve, calculated for the invariant melt in the volume (Mg) SiO - MgA^O^-KAlSiO^-KAlSi^Og ' or (^l/Sp/Kls/Lc) is shown on Figure 7. T^is

volume encompasses one of the largest compositional spaces in the system

and is therefore significant. The position of the equilibrium curve lies

close to the lower of those for the CaO-system.

References: J.F. Schairer, J. Amer . Ceram. Soc . 37./ 501-33 (1954)

W.C. Luth, J. Amer. Ceram. Soc. 50, 174 (1967).

d. Thermochemistry of the FeO^-K^O-Al^-SiC^ System (R.S. Roth, T. Megas

,

—— C. McDaniel, H. Parker,

H.I. Yoo, L. Cook and D.B. Minor)

To obtain a preliminary working model for this system under low

to moderate PQ pressure (prevalent in the higher temperature MHD regime,

where Fe > ^Fe+

^) , substitution of Fe for Mg is assumed thermodynamically

ideal

+^see Figure 8) so that subsolidus relations for Mg-phases also apply

to Fe -phases (this assumption breaks down in the SiO^-rich region) . Coal

ash analyses and tie lines are shown in the FeO^-A^O -SiO projections

of Figures 11 and 12. As noted above, FeO is an important constituent of

bituminous coals, comprising typically from 5-25 mole %. FeO in subbituminous coals is somewhat variable, with a maximum at 'v. 6 mole %. The

majority of analyses of both coal types project within the Fe

2 Si04

-KAlSi0 -KAlSi 0

system, for which data are lacking.

A thermodynamic analysis of the FeC>x-K

2 0-Al2 0

2 -Si0

2 system has been

begun as a prelude to study of invariant equilibria, according to methods

outlined a^ove for MgO- and CaO-systems. Calculation of the effect of

adding Fe to the volume 01/Sp/Kls/Lc indicates a rather marked lowering

of the curve in Figure 7.

The K O-FeO system is under detailed investigation. T. Takahashi,

K. Kuwabara and Kase [Denki Kagaku, 43_ (5) 273 (1975)] recently published

chemical and electrical data for this system. Our preliminary data are generally in accord with the above although melting temperatures appear

to be higher than reported. For example, the compound KFeC>

2 (hydrolyzes

at room temperature in air) is reported to melt congruently near 1325°C

whereas our experiments reveal melting above 1600°C. Most interesting is

the 6-alumina type phase in this system (solid solution between K 0*6Fe

2 0

2 and K 0*6.8Fe

2 0 ). We have identified the unknown compound reported at

K 0*10Fe

2 0 by the above workers as a 3" '-alumina type phase. The

electrical conductivity of these 3-alumina related phases is shown to

approach 1 (ohm-cm) ”1 at 100°C . These data prompted us to prepare several

ceramic specimens for preliminary measurements of electrical conductivity.

The base composition, K 0*7Fe

2 0 was selected either pure or doped with

Ti0

2 or Ta O.. Powders were prepared from K

2 C0

3 and Fe 2°3 wh^ ch were mixed,

calcined at 800°C and remixed. Discs were cold-pressed then isostatically

pressed and sintered at 1300°C covered with a powder of identical composition

(to minimize K^O loss)

The doped materials suffered from substantial

grain growth but remained suitable for preliminary measurements . High

36

37

Fig. 12. FeOxY-AloOo-I23 Si02 projection of analyses for bituminous coals

i|-fW

38

density bodies probably will have to be fashioned by hot-pressing.

Work is also being concentrated in the subsystems KFeO -SiO and

KAlSiO^-KFeSiO^ . The KFeO^-SiO^ system is analogous in mosi respects

to the KAlO^-SiO^ system (see previous Quarterly Reports) . Solid

solution of ^20 mol % SiO in KFeO^ stabilizes a cubic cristobalite-like

phase with a_=7.84A. A body centered tetragonal phase, K Fe Si 0

(a_=9 . 09 , £=5.33A), occurs at M3% si0

2

• KFeSiO develops at

1+X 1_X 4 low temperatures with the hexagonal kalsilite structure (£=5.278,

c_=8.824A) and transforms reversibly at 945 °C to a new orthorhombic

variant (a=9.155, b=5.433, £=59.61A). At about 25% KAlSiO a new

distorted form of kalsilite was found in the KAlSiO -KFeSiO system.

This form is monoclinic withaf=10 . 70 , b=8.55, £=10.65A and *8-120° .

e. Implications of Research on Seed-Slag Interactions (L. Cook)

The application of experimental studies to the understanding of

natural coal slags, especially the differences in behavior between

eastern and western coals, is clear from examination of Figures 4, 5, 6,

9, 10, 11 and 12. It should be emphasized, that when considering the

chemistry, data are most meaningful when treated on a mole % basis.

By studying melts of invariant composition within the significant

volumes of the multicomponent MHD slag system CaO-MgO-FeO^-K^O-Al^O^ -SiO^

,

it is expected that a maximum of information can be gathered per unit

time expended. This work will allow construction of a working model for

seed/slag interaction, once potassium concentrations have been measured,

related to subsolidus - derived potassium pressures, and converted to

potassium activity coefficients. This work will be closely correlated

with more detailed investigation of specific systems, and correlated when

possible with direct measurement of potassium vapor pressures. Most

important is the continuing investigation of the K^O-FeC^-A^O^-SiO^

system.

Presently available MHD channel slag analyses representing compositions

which appear to have been least contaminated by reaction with channel

components are given in Table 2. They contain from 15-25 mole % K^O.

These analyses are plotted in Figures 13 and 14 on (CaO, FeO )-Si0^-Al_0

3 "projections". All channel slags appear to differ from the parent coal

ash by having a significant enrichment in Al^o., an^ a Poss ikle slight

depletion in FeO . From these limited data ana the compositional analysis

given in FiguresX

4, 5, 6, 9, 10, 11 and 12, suggestions for the average

compositions of channel slag produced under MHD operation by "typical"

bituminous and sub-bituminous coals can be given, as in Table 3. It is

to be emphasized that these are suggested only as reasonable working

models and exceptions and local variations are expected to be the rule

(see Figure 15) . Particularly important are the effects of temperature

and P . Not all channel slag analyses and especially those in Table 2
'

.
39

CaO CaO

Fig. 13. Ca0-Al 2 03 -Si02 projection of channel slag analyses given in

Table 2

40

FeOx FgQx

Fig. 14. Fe0

x-Al

2 03 -Si02 projection of channel slag analyses given in

Table 2

,

41

0 400 800 1200 1600 2000 2400

Distance from plasma, jjm

Fig. 15. Variation in K2 O content across a specimen of channel slag from AVCO

Mark VI generator (S. Petty, pers. comm, to W. Hosier, 1975).

I

42

lend support to the idea that high CaO slags have a lower propensity for

K^O absorption. It may be that more data are needed before the effects

suggested by Cook e_t al_. (1976) to exist in synthetic compositions can be evaluated, especially for compositions with more than 30 mole %

(CaO + FeO + MgO)

An important point is that modeling of channel slag by systems with

KAlSiO^ and KAlSi^O^ as components (25.0 and 16.7 mole % K^O, respectively)

is required (Figure 6 and Table 2)

L.P. Cook, E.R. Plante, T. Negas, R. Roth, C. Olson, 1976,

Crystallization and Vaporization Studies on Synthetic

Coal Slag Compositions: Proc . 15th Symp. Eng. Aspects

MHD , Philadelphia.

Reference

43

Table 1. Gas-Solid Equilibria in the System CaO-K^O-Al^O^-SiC^

(apply also to invariant melts) •

Four Phase Volume Vaporization Equilibrium

1 Kls/Cor/CA /An

b

2Kls + CA = lAn + 6Cor + 2K + 1/2 0.

b 2 2-a Kls/CA /Ge/An

b 34Kls + 12Ge = 1CA. + 23An + 34K + 17/2 0.b 2 2-b Kls/CA^/ (Ge) 8 (KGe)

2/An 46Kls + 20 (Ge) (KGe) . = 1CA. + 35An + 50K + 25/2 0.

• 8 • 2 b 2 3-a Kls/Lc/Ge/An lKls + ILc + lGe = 2An + 2K + 1/2 0

2 3-b Kls/Lc/ (Ge) Q (KGe) ./An 12Kls + 6Lc + 10 (Ge) (KGe) = leAn + 20K +50

• 8 *2 2 4-a Kls/Lc/Ge/Wo 5Kls + 2Wo = lGe + 3Lc + 2K + 1/2 0

2 4-b Kls/Lc/ (Ge) (KGe) _/Wo

• © • 2 42Kls + 18Wo = 10 (Ge) _ (KGe) . + 24Lc + 16K +40 «o«2 2

5 Kls/Lc/Cor/An 4Kls = 2Lc + ICor + 2K + 1/2 0

2

6 Lc/Cor/Mu/An 2Lc + 5Cor = 2Mu + 2K + 1/2 0

2 7-a Lc/Ge/An/Wo 3Ge + 4Lc = 5An + lWo + 4K + 0

2 7-b Lc/ (Ge) (KGe) ./An/Wo

• o • 2 30 (Ge)

8 (KGe ) 2 + 30Lc = 42An + 12Wo + 36K + 9 0

2

8 Lc/Ksp/An/Wo lWo + 5Lc = 3 Ksp + lAn + 2K + 1/2 0

2

9 Lc/Ksp/Mu/An 16Lc = IMu + 10 Ksp + 6K + 3/2 0

2

10 Ksp/Mu/An/Q 6Ksp = IMu + 16Q + 6K + 3/2 0

2

11 Ksp/An/Wo/Q lWo + 2Ksp = 5Q + 2K + 1/2 0

2

Kls = KAlSiO . = kalsilite

4 Lc *; KAlSi.O. = leucite

Ksp = KAlSi 0 = potassium feldspar

Cor = a^2°3 = corun<*um • Mu 3Al

203 *2Si0

2 = mullite

<*6 - CaA112°19

\

Ge - Ca

2Al

2 Si0

7 = gehlenite

(Ge) 8(KGe)

2 = (Ca

2Al2 SiO

7

) 0>8 (KCaAlSi

2O7

} 0.2

An “ CaAl

2Si

2 Og = anorthite

Wo * CaSiOg = wollastonite

Q

« Si0

2

*= quartz

Table 2

Channel Slag Analyses (Recalculated on basis of 100 Mole % - SiO„+Al 0 +FeO +CaO+MgO

2 2 3 x or Si0

2+Al

20

3 +Fe0x+Ca0+Mg0+K2 0 )

(Si0

2+Al203

) .7734 .7317 .7169

A. AVCO Mark VI Test 8/6/74 1

1 Parent

Coal

Ash

2 Burner

. Slag

(-k

2o)

3 Channel

Slag

(-k

20)

Burner

Slag

(+k

2 o)

Channel

Slag

(+k

2 o)

Si0

2 61.47 59.57 58.24 49.19 48.49

W2°3 18.01 21.84 23.00 18.04 19.15

FeO

X 16.69 14.67 14.03 12.12 11.68

CaO 2.52 2.33 3.16 1.93 2.63

MgO 1.31 1 .59 1.57 1.31 1.31

V

SiO

17.43 16.75

1 AVCO Everett Res . Lab . , Experimental MHD Powder Plant Development

Report, October 1, 1974.

Program, Monthly Progress

B. AVCO Mark VI Test 9/20/74 2

1 4 5 Parent Cathode Anode Cathode Anode

Coal Slag Slag Slag Slag

Ash (-k

2o) (-k2o) (+k20) <+k

2

o)

Si0

2 61.47 57.82 55.17 43.27 45.49

ai

2o

3 18.01 25.17 29.13 18.84 24.02

FeO

X

16.69 12.10 11.28 9.06 9.30

CaO 2.52 3.43 2.99 2.57 2.47

MgO 1.31 1.47 1.43 1.10 1.18

V — — — 25.17 17.54

SiO,

.7734 .6967 .6544

(Si0

2+Al

203

)

AVCO Everett Res. Lab., Development Program for MHD Power Generation, Monthly Progress

Report, November 1, 1974 • \

C. AVCO Mark VI B Test 4/29/75 3

Cathode Anode Cathode Anode

Slag Slag Slag Slag

(-*

2o) (-k

2o) (+k

2o) <+k2o)

Si0

2 64.82 63.47 52.00 53.79

u20

3 23.59 25.99 18.92 22.03

FeO

x 6.58 4.77 5.28 4.04

CaO 3.28 3.89 2.63 3.30

MgO 1.73 1.88 1.39 1.59

K,0 — — 19.78 15.24

^ AVCO Everett Res. Lab., Development Program for MHD Powder Generation, Progress Report

Report No. 46, July 1, 1975.

D. Stanford Facility, Teat #9, 19764

6 7 Parent Channel Channel Coal Slag Slag Ash <-k

2o) (+K

20)

S10

2 67.19 55.08 46.72 m2o

3 23.59 29.61 24.75 FeO

X 5.51 3.67 3.07 CaO 2.33 1.31 1.09 MgO 1.39 9.53 7.97

*a° — — 16.40 SiO,

<SiO

2+Al203

) .7401 .6536

8 . Stanford Facility, Test #11, 1976'

SiO„

<Si02+Al

203

)

8 Parent Coal Ash

9 Channel

Slag

(-k

2o>

.3088 .7419

Channel

Slag

(+k

2o)

sio

2 56.73 54.69 45.96 A!

203 13.41 19.03 15.99 FeOX

4.42 4.28 3.60 CaO 18.89 8.78 7.38 MgO 6.55 13.22 11.11 *2° — — 15.96

J.K. Koester, M.E. Rodgers, and K.R. Eustis, 1976, In Channel Observations on Coal Slags Proc. 15th Symp. Eng. Aspects MHD, Philadelphia.

r. Stanford Facility, Test #12, 1977 5

8 Parent

Coal Ash

sio

2 56.80 ai

2o3 13.42 FeOX

4.38 CaO 18.87 MgO 6.52 V —

SlOj

(S±0

2+A1

20

3

) .8089

10 Channel

Slag

(-k

2o)

Channel

Slag

(+k

2o>

50.23 42.58

21.03 17.83

3.38 2.87 18.27 15.49 7.09 6.01 — 15.22

.7049

G. Stanford Facility, Test #19, 1977' X

11 12 Parent Channel Channel Coal Slag Slag Ash (-KjO) (+k

2o)

sio

2 67.23 71.51 59.68 ai

2o

3 13.69 19.94 16.64 FeO

X

12.16 7.97 6.65 CaO 6.92 0.59 0.49 MgO — —

*2° — — 16.54

SiO.

.8308 .7820

[Si0

2+Al

203

)

8 J.K. Koester and R.H. Nelson, 1977, Electrical Behavior of Slag Coatinga in Coal-Fired HKD Generators: Proc. 16th Symp. Eng. Aspects KHD, Pittsburg.

Humber* above the analyses refer to data points in illustration*
V

'
46

Table 3. Suggested Average Compositions for Modeling MHD

Channel Slag (Mole %) .

Western Coal Eastern Coal

Si0

2 42.0 46.0

A1 2°3 17.5 19.0

FeO

X

5.0 12.0

CaO 12.0 2.4

MgO 6.0 1.2

K2° 17.5 19.4

100.0 100.0

sio

2 Si02+Al

2 0

3

0.705 0.705

47

Diffusion in Insulator-Electrode Couples (A. J. Armstrong, E. N. Farahaugh

J. R. Manning)

Insulator-electrode couples, as described in the March 1977 report,

have been made. These couples were heated at 1600 °C for 1, 10 and 100 hr.

periods. Preliminary Fe diffusion data has been taken on a MAFF-31 - MgO

couple

Figs. 16, 17 and 18 are optical micrographs of MAFF-MgO couples heated

for 100, 10 and 1 hour respectively. The 100 hr. test drove enough iron

from the MAFF into MgO to cause a deep rust-colored discoloration of the

whole of the MgO insulator. The 10 hr. test also produced a discolored

band in the MgO as can be seen at the edge of the MgO (middle layer) in Fig. 17

The 1 hr. test produced a much narrower band in the MgO as can be seen by

comparing Fig. 17 and 18.

In these couples the boundary between the insulator and electrode was marked by fine Pt wire placed between the two materials. Fig. 19 is a 100X SEM micrograph revealing that the MgO and MAFF were actually held

apart ^25 ym by the Pt wire. The wire is seen as the white appearing

round spot about midway along the boundary of MgO (the coarse grained

material) and the MAFF-31.

The Fe concentration profile in the MAFF and MgO as determined by SEM

and EDX techniques is shown in Fig. 20. The boundary between the two

materials is indicated by the long line. Note the decrease of Fe towards the

edge of the MAFF as expected. We also see a hump in the Fe profile in the

MgO about 75 ym from the boundary. In previous measurements on MgAl 2 0^-MAFF

diffusion couples, we have also seen and reported such a hump in the Fe

profile, but it was reported to lie in the electrode material. In Question

in those tests was the exact position of the electrode insulator interface.

It was to locate the interface that the Pt wire markers were introduced.

On the basis of this preliminary data, it seems that the hump lies

in the insulator material for this MgO-MAFF couple. Also, it appears that

the fact that the electrode and insulator were not in intimate contact, did

not inhibit diffusion of Fe from the MAFF into the MgO.

Profiles will be taken from the other couples and plotted. Particular

attention will be paid to the position of any Fe hump relative to the

insulator-electrode boundary.

1
:
48

Fig. 17

Optical micrograph of MAFF-31-MgO couple tested for 10 hrs.

49

Fig. 19 . 100 SEM micrograph showing separation of MgO (coarse-ground material)

and MAFF-31.

50 j

51

Task K. Materials Testing and Characterization

1. Introduction

During this reporting period an appreciable fraction of our efforts

have been devoted to characterization of materials, design and fabrication

of electrode/insulator modules, and evaluation of tested electrode assemblies.

This includes structural analyses of ceramic powders and solid samples

from various sources (sections 2a and 2b) consultations and preparations

for future tests (section 3) , results of the analyses on 3 electrode modules

(TT/NBS, ANL and GE) tested in the MIT-rig (section 4a,b,c), as well as examinations of Fluidyne preheater materials (section 5a) and fly-ash

(section 5b)

2. Structural Analysis of Powdered and Solid Ceramics

a. X-ray Diffraction of MHD Materials (C. L. McDaniel)

During this quarter we received over 35 materials (powders, dense ceramics,

etc.) from various external sources. Most of these were fabricated ceramics,

sinterable or sprayable powders or powders at various stages of processing.

These were analyzed by powder x-ray diffraction methods. Data, transmitted to

the appropriate source, are shown in Table 1 (following p. 52). Secondary phases

(e.g.

,

at grain boundaries, segregations, etc.) when present (depending on

the phase and host material) at levels below roughly 3% are not detected by

routine x-ray diffraction. If requested, samples will also be analyzed for

minor phases using SEM/EDX methods.

b. ANL U-02 Phase III Materials (T. Negas, C. L. McDaniel)

ANL delivered five materials related to their U-02, Phase III effort.

X-ray diffraction analyses were requested. Treatment of these materials

was fully documented by ANL and is summarized together with x-ray results

below.

MHD #1 - 18 m/o Ce0

2 -82Hf0

2 MHD# 2 -4 - 9.7 m/o Y 0 • 16 . 3CeC>

2

• 74HfC>

2 Powder - 18 m/o CeC>

2 -82Hf0

2

, commercial
'

'
52

SAMPLE TREATMENT X-RAY RESULTS

Powder none Roughly equal proportions of

monoclinic Hf0

2 ~type solid

solution plus a cubic or tetragonal phase rich in HfO . Broad x-ray lines suggest fine

grain size and/or poor crystallinity

MHD #1 hot-pressed, vacuum ^1500°C,

5600 psi, ^40 min. Annealed

in air 1000-1150°C

90% dense

Monoclinic Hf0„-type solid solution

with a trace of CeO„

MHD in hot-pressed, vacuum,

VL400-1435 0

, 5000 psi,

^20 min. Annealed

1150°C in air. 95% dense

Cubic fluorite solid solution

(Hf0

2 ~rich) of excellent crystallinity (aj=5.144A) plus monoclinic

HfO (<5%)

MHD #3 hot-pressed, vacuum

^1300-1325°C, 5000 psi,

^20 min . No anneal in air

95% dense

Same as MHD #2 but with slightly

more monoclinic HfO_. Cubic phase

has a=5.153A suggesting a somewhat

lower Hf0

2 content than in MHD #2.

MHD #4 hot-pressed, flowing Ar,

1300-1325°C , 5000 psi,

n-15 min. No anneal in air.

82% dense.

Cubic fluorite plus monoclinic HfO,

(in roughly a 2:1 proportion) plus'

a trace of Y-0-. Not as well

reacted as MHD J

//2 and #3.

3. Preparation for Future Tests (W. R. Hosier)

A considerable amount of time has been spent in consultation with

Westinghouse on planning the materials aspects of Phase III of U-02 . In

particular, there has been close contact with APS Materials, Inc., Dayton,

OH and Trans Tech, Inc., Gaithersburg, MD in an effort to obtain the best

possible powders and sprayed material for certain electrodes in the Phase III

design. It is necessary to maximize the density and obtain single phase

pore-free material to reduce MAFF-31 oxidation and seed penetration. At the same time, resistance to fracture must be high. In arc plasma sprayed

LaCr03 (Mg) the density and resistance to fracture must also be maximized along

with obtaining single phase pore-free material. A program has been initiated

with APS Materials, using arc plasma sprayed powders from Trans Tech and others

to study these requirements and develop these materials.

Several proof tests for Phase III U-02 are scheduled for August and

September at Westinghouse R & D Lab. , Pittsburgh, PA. Subsequent analysis

of the electrodes will be made and the results will be used to determine

the materials and electrode designs to be used for the Phase III test

scheduled in Moscow in early 1978.

53

Table 1. spravable powders

Comments on

aterial From Supplier Description X-ray Diffraction Availability

«"5 .05CrO

3 NBS 6500 A-1A

CERAC Crushed aps powder

-200 +25u mesh 1% excess Cr^O^ added

single phase

LaCrO^-crystallinity fair to good

Special order

2-3 weeks delivery

any mesh size

Trans-Tech spray dried aps powder

6-8 weeks delivery

special order, any mesh size

| 3 Westinghouse

(Kochka)

LC-GR02

Gen. Refract. no details supplied,

aps powder crushed

med. good crystallization, no para- mater shifts, no 2nd phases

t°3 Westinghouse

(Kochka)

LC-CEOl

CERAC aps powder, crushed

no further details

given

med. crystallization, no parameter

shifts, no other phases

^2°4

- Fe 3°4 NBS Trans-Tech aps powder spray dried -140 +325

single phasegood crystallizationOff shelf item

*°3 Westinghouse

LC-CE02

CERAC fines from 6500

A-1A

single phase, good

crystallinity

2 to 3 weeks del.

•1 2°

4 NBS S-71 Trans- Tech

Type 1 aps spray dried

powder, -140 +325

single phase, good

crystallinity

Off the shelf

V4 Westinghouse

(Kuszyc)

MA-CE02

Not stated aps powder crushed

98% pore, -150 +325

mesh

single phase

good crystallinity

:rC3 NBS Lot.

No. 4283

CERAC aps powder crushed,

-150 +325 mesh

cubic perovskite

single phase

Off shelf-stock * 1123

W3

|

NBS Lot.

No. 3451 B CERAC aps powder crushed,

-325 mesh

cubic perovskite

single phase

Dff' shelf-stock # 1129

f°3

Westinghouse

(Kuszyc)

SZ CEO

2

No origin

given

aps powder, -150

+325 mesh

single phase , good

crystallinity

:r°3

J

Westinghouse

(Kuszyc)

SZ DM02

no origin

given

powder-no details

given

single phase, poor

crystallinity

_

:r°

3

I

NBS Lot

No. 3757

CERAC aps crushed powder

-150 +325 mesh

orthorhombic perovskite + second phase

of 10-20% cubic Ca stabilized ZrO^

Off shelf-Stock Nc. 1021

*°3

||

NBS Lot

No. 375

CERAC aps crushed powder,

-325 mesh

orthorhombic

perovskite plus

5 to 10% Ca stabilized Zr0

2

Off shelf -Stock No. 1022

V12Y2°3 NBS 202NS

METCO aps power mechanical mix of Zr°2~Y 2°3

Off Shelf

k NBS CERAC aps powder crushed

-150 +325

single phase CeO^

good crystallinity

Off shelf - Stock No. 1232

3

3 NBS CERAC aps powder crushed

-150 +325

single phase Y^Ofair crystallinity

Off shelf - Stock No, 1584

V5T2°3 NBS CERAC aps powder crushed

-150 +325 mesh

cubic HfO_ plus

5-10% monoclinic

HfO^ - crystallinity

very poor

Off shelf - Stock No. 1547

Oj- 12 m/o

Oj - 3 m/o

1°3

Westinghouse

(Kuszyc)

CZ-CEOl

no origin

given

aps powder, no mesh

given

broad peaks, mostly

cubic - small amount tetragonal poor

crystallinity
!

•

'

,
54

LaCrO

3 Westinghouse

(Kochka)

LC-GR05

Gen. Refract.

SrZr0

3 Westinghouse

(Kochka) ZA

no powder

source given

SrZr0

3 Westinghouse

(Kuszyk)

3A-SZ UM02

no powder

source given

Ce0

2 -Y

2 03 -Zr0

2 onCa(Mg)Cr0

3

Westinghouse

(Kuszyk)

no source given

MgO, AI2O4

(MgAl2 04

) Westinghouse

(Rossing)

MA-CEO301

CERAC

SrZr0

3 Westinghouse

(Rossing)

SZ-CE0301

CERAC

La

.95M9.05 Cr°3 Westinghouse

(Rossing) CERAC

sinterable powder good crystallization

no parameter shifts

no 2nd phases

sintered 1650 °C no time, atmosphere

given

single phase

crystallinity very good

sintered 1650 °C no further datasingle phase

crystallinity very

good

flame sprayed

piece, no proportions given for

Ce02-Y

203-Zr°2

top surface Ce02~Y2(-> 2~'z‘ z:<-> 2 broad peaks perhaps single

phase cubic type, possibly

some tetragonal ZrC>2 type

lower -LaCrO -poor crystallinity +2 to 5% unknown phase

as in LC CE 0201

plasma sprayed

piece made from Cerac Mg0-Al o 0

3 powder -325

+10u mesh

single phase crystallinity

medium to good

plasma sprayed

piece made from Cerac SrZr0

3 powder -325 +10u mesh

SrZrO +5-10% tetragonal ZrO^

crystallinity SrZr0

3 good

ZrC>

2 poor

plasma sprayed

piece made from Cerac powder

LC-CE0201

LaCrO +unknown+La

2 03 unknown phase 10% La2°2 5% crystallinity poor

£

55

The electrodes (22) for the AVCO test in a slagging environment are in

various stages of fabrication. Six electrode pairs have been completed by

APS Materials, Dayton, OH. Five more electrode pairs are in progress at

Technetics, Division of Brunswick, Milford, CT and should be completed soon

No date has been set for the test run, but it should occur during the next

reporting period. A complete description of the design and test configuration will be given later.

4a. Examination of NBS LaCrO

^ Electrodes (E. N. Farabaugh)

Anode and cathode material of cermet bonded LaCrC>3 with a Zr02 cap

has been examined by SEM and EDX techniques. These electrodes were the

subject of test #166 on the MIT test rig on 2-18-77. The test ran for

3 hr. and 52 min. at temperatures of 1700 °C. The electrodes were under

power and seed conditions for 2 hr. and 41 min. of the test. A drawing

of the electrode assembly and cooling strip is shown in Fig. 1.

A cathode, after potting in epoxy, was cut along its length to

provide a cross section of various positions along the long direction of

the electrode. Areas where the Zr02 cap plus a little LaCrO^ had come free from the main LaCrC>3 body as well as areas where the cap was well

attached to the LaCrO were observed. Generally, the bond between the

cap and the LaCrC>3 had held up well under testing. However, the entire

cathode assembly had come free from the copper cooling strip leaving a thin layer of the Cu-Al alloy paste, used to bond the electrode to the

cooling strip, on the bottom of the electrode.

Fig. 2a is an optical micrograph showing a typical cross section

near the center of the electrode. The Zr02 (^.78 mm) is the top white

layer. Next is the LaCrC>3 (^2.3 mm) and at the bottom is the Cu-Al

alloy. There is no evidence of the cap coming loose in this area.

Fig. 2b shows an area where the cap has come free from the LaCrC>3

. This is from a part of the electrode next to that shown in Fig. 2a.

It is noted that there is some LaCr03 with the cap suggesting the failure

came in the LaCr03 instead of at the Zr02 -LaCr03 bond. Again, the Cu-Al

alloy is seen at the bottom indicating the apparently poor bond of the

paste to the cooling strip.

Fig. 2c is a 20X SEM picture of the cathode showing the separation of

the cap and main body. There are several interesting areas to discuss.

The first, A, is the cap and a little of the LaCr03 which has separated.

There is no obvious definition of the boundary between the Zr0 2 and LaCrC> 3 , indicating a good bond. Some K was detected in this area using EDX methods

B is a narrow layer in the LaCr0 3 which has a K concentration not much

different from that in layer A. Area C is a layer in the LaCrC>3 which has

a much higher concentration of K than either the material just above it (B) or just below it (D) . The material in C also shows a much different microstructure than either B or D. D is the last of the LaCr02 before the Cu-Al

alloy layer, layer E. K has been detected throughout the entire assembly.

Some K was found even in the Cu-Al alloy layer.
'

.

-
56

Close examination of EDX spectra showed a decrease in the Cr peak in

layer C relative to layer B. This could be explained by the K reacting

with the Cr and reducing the presence of Cr in the LaCrO^. An exact

analysis of the spectra is complicated by the overlapping of CrK,^ and LaLg

lines, and other elements (i.e., scandium). In layer C it seems reasonably safe to say that higher K concentrations are accompanied by

lower Cr concentrations.

It was noted that the separation of the cap occurred above where a well defined, K-rich layer was formed in the main body. It was further

determined that the place where the Zr02 cap and LaCrO^ broke free from

the rest of the LaCr03 appeared to be where the Zr concentration has dropped

to near zero.

We must consider two results: (1) a K-rich LaCrC>3 layer is found

where the cap has parted; ( 2 ) the separation has occurred where the graded

ZrC>2 has little Zr concentration. One cannot say which has occurred first.

It may be that the separation started because of some thermo-mechanical

strain, allowing K to easily get to the LaCr03 through the porous Zr02

causing expansion which caused the cap to break away carrying part of the

LaCrOj with it.

The top part of the electrode was exposed to temperature from 1300 °C

to 1700 °C at the top of the cap. At these temperatures, K will not condense,

hence, the small concentration of K in A and B. Layer C could correspond

to the temperature range of 1100 ° - 1300 °C where K is a liquid. (K2 CO 3

) The band

C had the largest concentration of K and the most microstructural changes

which might be expected if the K-salt was liquid in that region.

Anode material, after potting, was cut to give a cross section at right

angles to its length. Optical examination of the anodes showed that in none of them did the cap come free from the main LaCr03 body. However, some Cu-Al

alloy was still on the bottom edge of the electrode showing that the anodes,

as well as the cathodes, came free from the copper cooling strip at the alloy

paste-copper bond.

EDX examination of the anode indicated the presence of K, Si, Mg, Zr

,

La, Cr, Y and Sc. There was not as much K present as in layers A and B of

the cathode.
“l

'

*

.
APS

MqAUOy.

INSULATOR

57

cr

LlI EE

ro

CM

ro Ooo

_]aQO

GO

hLU

CC

LUO

CD

cntn0) X!0G

-Po0

U-lo

01£

:s(0pQ

Oi

•H

fa

J

58

59

4b. Post Test Analysis of ANL Electrode Test at MIT (T. Negas)

On March 18 we received four electrode assemblies, including virgin

specimens from ANL. These were tested in the MIT facility, March 8 . The test was terminated after 30 minutes and no seed (or current) was utilized. Only one electrode remained intact partly as a result of

difficulties related to the set-up and the running of the test as well

as the incompatibility of the ANL electrode design (originally for the

Reynold's facility) with the MIT test facility. These considerations

together with observations during the test are documented by an ANL

report, "Synopsis of ANL Electrode Test at MIT 8 March 1977."

The electrodes had a configuration illustrated in Figure 3 (provided

by ANL). The ceramic portion consists of 82Hf02 , 8 Y 2 O 3 and 10CeC>2 (HYC) . Three of the electrodes incorporated a mesh of Nb (1% Zr) wire (apparently

clad with an Au/Pt combination) while the remainder contained nichrome wire.

Only the latter survived the test basically intact. Figure 3 suggests that

the insulation is A^O^ (plasma sprayed) . X-ray diffraction and SEM/EDX

data indicate that this material is actually MgAl 204 (spinel) . Zirconia

castable was used to fill gaps between electrodes when stacked.

Figure 3. Schematic configuration of ANL electrodes (provided by ANL).
'

.

-
60

A separate report (11 figures) of our analysis was submitted to

ANL . Conclusions based on this report are summarized below.

la. The basic ANL electrode appears encouraging from a thermal

design standpoint. Combining wire metal within an integral

cohesive electrode unit is a sound approach for two reasons:

1. the wire, well into the ceramic, provides for

electrical conductivity and,

2 . means of attachment to an underlying metal is

provided.

lb. The design was not subjected to potassium and E field,

conditions which may alter the electrode structure.

2 . The ceramic oxide-nichrome electrode survived the test

in relatively good condition (see Figure 4)

3. The ceramic oxide-Nb (1% Zr) wire electrode (3 of 4)

were damaged extensively (see Figure 5)

a. The Nb wire oxidized and reacted with adjacent

ceramic components. Localized melting followed

and some liquid was injected into the side

insulation

b. "Hot spots" (>1900°C) observed during the test

contributed to localized heating and oxidation.

4. Fissures (and detachment) are apparent at the ceramic-braze

interface

5. The plasma sprayed MgAl 0^ insulation appears to bond well

to the ceramic oxide but poorly to metallic components

6. The ceramic oxide as fabricated contains coarse grains

(HfC^ , monoclinic) surrounded by a finer matrix (CeO , Y^O^-rich) . This matrix may provide avenues for seed

penetration and subsequent corrosion of construction

materials
'
61

Fig. 5 . Post-test ANL electrode K9X) fabricated from ceramix oxide (see

text) plus Nb (1% Zr) wire. Oxidation of the wire is evident above

the dashed line. Z = zirconia castable; i = MgAl^O^ insulator;

M = several areas wherein chemical interaction including melting

have progressed.

Fig. 4 . Post-test ANL electrode (^9X) fabricated from ceramic oxide (see

text) plus nichrome wire. Z = zirconia castable; i = MgAl^O^ insulator.
.‘

'
62

4c . Post-test Analysis of G.E. Electrode Assembly Tested in MIT Rig . (L. Cook)

A cathode assembly (#154-CA) and anode assembly (#154-A2) from a

4 hr test performed by M.I.T. on January 13, 1977 were provided by the

General Electric Company for materials characterization. Potassium

carbonate was the seed material and no slag was introduced. Other details

of the test and unused specimens for comparison were not available. The

electrodes are of a "proprietary design." For purposes of characterization

the essentials are: 1) (Mg ,A1 ,Fe) - spinel of 2 mm thickness bonded to

2) Mo-sheathed, finned copper lead-out, in turn bonded to 3) a nickel

base with 4) (Mg, Al) - spinel slabs as insulators between electrodes.

The entire electrode assembly is attached in a complicated way to a copper

cooling block.

The anode and cathode assemblies appear to have behaved similarly in

the 4 hr test, except: the anode may have been slightly more resistant to

corrosion; the anode insulator underwent markedly stronger attack than the

cathode insulator. Whether these differences are significant or are a peculiarity of sampling is unknown. The spinel insulator material appears

to have withstood the test well (Figure 6) , with the exception of local

breakdown in the anode insulator.

It is estimated that a minimum of 25% of the spinel electrode material

was removed during the test. In some areas 100% has been removed. On the

average probably 50% was removed. Once degradation to the level of the

lead-out assembly occurred, corrosion of the lead-out fins, possibly by

local melting, was rapid in some areas. The ceramic-metal bond at the

spinel/lead-out junction appears to have survived the test.

The mechanism of degradation of the spinel electrodes may be largely

thermo-mechanical. This suggestion cannot be evaluated without knowledge

of the thermal history, but is given support by the observation of partly

developed cracks which do not appear to be due, for example, to any chemical

process related to penetration of seed. They do appear related to the

highly porous areas in the spinel. To what extent textures in the spinel

(Fig. 7) are a function of fabrication vs. MHD operation cannot be evaluated

without comparison with fresh electrodes.

*Substantial portions of this report have been deleted, as they divulge

details of a proprietary nature. The material presented is done so with

the permission of Dr. M. Noone , General Electric Corporation.
.

-
63

Fig. 6. Secondary electron image of insulator surface, G.E. cathode

assembly (800X) , Plasma flow was from top to bottom.





66

5a. Examination of Fluidyne RFG Samples (A. Perloff and E. N. Farabaugh)

Three samples of a chrome-magnesia refractory material were examined by

x-ray powder diffraction techniques. The samples were (1) original

(untested) material, (2) FN-96 cycled for 30 hours between 1800-1600°C

with ash but no seed and, (3) FN-91 cycled between 1850-1700 °C with ash

and K2 SO4 seed. For the tested samples the material to be examined was chipped from the reacted surface layer.

Original Material :

O The sample is predominantly well-crystallized MgO (a = 4.212A) and

a single phase spinel (a = 8.278A). This spinel is very close in cell

size to that reported for magnesiochromite [(Mg, Fe) (Cr , Al) 2 O 4 ] , but the

intensities do not correspond very well. This suggests that the cation

distribution may be a variation of the magnesiochromite. There are traces of materials which are reminisci.ent of forsterite

(Mg 2 Si04 ) and monticellite (CaMgSiO^) but the fit is far from perfect

and the identification must be considered tentative.

A sample of untested rebonded chrome-magnesia refractory (RFG) was characterized by SEM and EDX techniques . SEM micrographs showed the

material is to be composed of many large grains. EDX spectra from these

large grains revealed the presence of Mg, Al , Si, Ca, Cr and Fe , This

suggests the presence of several phases in the grains. Further examinations

of the large grains showed well-defined smaller grains of phases rich in

Mg, Al, and Cr suggesting a chrome-spinel phase. X-ray analysis also indicated

Si and Mg concentrated in areas surrounding the chrome-spinel phase. Thus,

it is possible that MgSiC>4 is present between the chrome-spinel grains.

MgO was easily seen in the structure when Mg mapping was done.

The untested structure is typified by large grains composed of several

phases. MgSiO^, MgO and chrome-spinel were predominately seen.

Tested RFG will be examined next and EDX characterizations will be

used together with X-ray powder data to better describe the phase structure.

FN-96

O This sample consists primarily of MgO (a = 4.218A) and two spinel phases.

The proportion of MgO relative to the spinel seems to be somewhat smaller

than in the untested material, but the difference is small enough that it

could be accounted for by intensity variations due to differences in

crystallinity (likely) or preferred orientation (unlikely) . The increase

in cell size of the MgO relative to the other samples is small but real.

The larger cell in this sample suggests some substitution of larger cations,

probably Ca and/or Fe, which has not occurred in the other tested sample.
.

.
67

The two spinel phases have much broader peaks than the untested

material spinel. This could indicate a much smaller crystallite size,

a lower degree of crystallinity or possibly each phase represents a narrow range of cell sizes clustered around the average values of

8.25& and 8.19iL Both cell sizes are smaller than the cell for the

original material which indicates the loss of some of the larger cations

(Fe and Cr) . There is more of the 8.25A material than of the 8.19 A phase

(contrast with results on FN-91) . This may indicate a smaller loss of Fe

and Cr than in the FN-91 sample. However, it could also be due to the

greater thickness of this sample providing proportionately more lower

temperature material.

There is significantly more forsterite (Mg 2 Si04 ) in FN-96 than in

the other two samples. In addition, there are two weak lines (d = 2.55,

3.21) not explainable by any of the above phases. These lines could be

due to an enstatite-like phase (MgSiC^) which contains some Fe substitution for the Mg, but this identification is highly uncertain.

FN-91

0 This sample is predominately a two phase spinel system (a-^ °^8.23A, a2 ^8.18ji) plus some MgO with traces of ether material. The most striking

difference relative to the other samples is the reduced amount of MgO

compared to the amount of spinel. This difference must be due to reaction

with seed material. The MgO that is present has the same cell as pure

MgO which indicates no substitution of Fe or Ca has taken place in

contrast with the FN-96 sample.

The spinel peaks are broad as in the FN-96 sample which suggests the same possibilities with regard to crystallite size, crystallinity and

variation of cell sizes. The spinel cells are similar or slightly smaller

than those found in FN-96. In the FN-91 there is more of the 8.18A material

than of the 8.23A phase. This could indicate a greater loss of Fe and Cr

due to the presence of seed or the slightly higher temperatures. However,

this sample was not as thick as for FN-96 so this difference could be simply

indicative of a greater proportion of surface material in the sample.

The impurity levels are too low to make identification certain, but

there is probably forsterite (Mg 2 Si04 ) and moniticellite-like (CaMgSiC^)

material. Additional possibilities include traces of KAlSiC>4 and

potassium magnesium silicates.
.

J

-

,
68

5b. X-ray results on slags from AVCO generator using Montana Fly Ash and Mew

Hampshire Fly Ash slags (A. Perloff)

We received from AVCO specimens of generator slag originating from

Montana and New Hampshire fly ash. It was requested that these be

analyzed for phase distributions, segregations, inhomogeneities, etc.

which could account for observed polarization phenomena during testing.

Our first step has been to identify the crystalline phase assemblage

of each slag specimen in bulk (see below) . SEM/EDX analysis of these

materials have not been concluded.

Six slag samples were examined by x-ray diffraction techniques. They

are derived from two sources of fly ash and were identified as New Hampshire

fly ash (anode and cathode) and Montana fly ash (front anode and cathode

rear anode and cathode)

ANODES

There is more general similarity between the N.H. fly ash pattern and the

Montana fly-ash from the front electrode. Both are primarily orthorhombic

KAlSi04 patterns. Impurity lines are d = 3.73, 3.27, 3.00, 2.90, 2.88, 2.72,

2.53 (N.H.), 2.38 and 1.936. The Montana fly-ash from the rear electrode

contains orthorhombic KAlSiO^, but the intensities are not normal for many

of the lines which could be due to more substitution of other elements or different K, Al, silicate phases (hexagonal and tetragonal phases are possibilities) . There are some lines appearing which are not significant

in the other anode patterns, e.g., (d = 4.46, 2.83, 2.75, 2.66, 2.59, 1.912,

and 1. 880) .

The complexity of the patterns make it impossible to unambiguously

identify the impurity phases. A number of silicates of Ca, Mg, Al , and

K are possible contributors. The probable presence of the tetragonal

potassium aluminum silicate in the Montana fly ash rear anode suggests

this may be richer in potassium than the others (lower temperature?,

greater accumulation in rear of channel?)

The only anode sample which could contain significant amounts of

iron oxide is the New Hampshire fly ash (d = 2.53).

CATHODE

All three cathode slags contain orthorhombic KAlSi04 but in reduced

amounts relative to the anode slags. The cathode slags seem to be

richer in the cubic and tetragonal phases which are higher in potassium.

This is more pronounced in the Montana slags than in the New Hampshire

one. The New Hampshire slag, although containing less of the potassium

rich aluminosilicates is the only pattern to show a clear indication of

KHCO^ lines. No iron oxides or 8-alumina phases appear to be present.
/
69

Task L. Assessment of Steam Plant Components (J . R. Cuthill)

This is the 8th quarterly progress report since the initiation of

Task L. The overall objective of Task L is the recommendation of

commerical alloys for critical components downstream of the MHD

generator. The assessment of these commerical alloys is being made on the basis of data in the literature, including research papers,

handbooks, data compilations, producers bulletins, specifications, and

the ASME Boiler and Pressure Vessel Code. If specific construction

alloys for these critical components cannot be recommended per se on the

basis of the data in the literature, promising alloys will be cited and

recommendations made for further testing. Such testing will be carried

out under this Task to the extent possible.

Table 1 is a compilation of mechanical properties of some promising wrought alloys with the steam superheater tube application

principally in mind. Table 1 is an extension of the corresponding Table 2 in

the previous quarterly report. (Table 2 in the present report gives the

corresponding alloy compositions). Each of these alloys has been

included because of relevance to downstream component applications, i.e.

simulation corrosion test data is available on them, or they are commonly used construction materials for steam plant components, or on the basis of related performance they appear to have promise of

withstanding the unique service conditions imposed by the presence of

the potassium compounds, which has been explored in previous quarterly

reports. Also, in the quarterly report for Jan. 1 to Mar. 31, 1977, a relative alloy ranking is given in respect to anticipated ability to

resist the hot corrosion attack. Conversion factors are given as a footnote to Table 1 for converting the mechanical property values in

the Table, given in SI units, to the English units which continue to be

used in the ASME Boiler and Pressure Vessel Code. (1)

ASME Boiler and Pressure Vessel Code

Many of the alloys in Tables 1 and 2 have been tabulated in the

ASME Boiler and Pressure Vessel Code. These alloys are listed in Table 3 together with their respective ASME Specification No. After each alloy

there is listed for comparison the maximum tensile stress allowed by the

Boiler Code at 1200°F, the estimated external surface temperature of the

steam superheater tubes. The maximum temperature at which the alloy can be used is also given together with the corresponding maximum allowable

design stress. Note that in some cases this maximum allowable service

temperature is below 1200°F (648 C)

Hard Facing and Wear Resistant/Corrosion Resistant Alloys

There are applications such as the mating surfaces of the valves

on the high temperature air preheaters, that will require alloys to

withstand metal- to-metal sliding wear as well as high temperature

oxidation and hot corrosion. Some pertinent properties of two selected

70

alloys are given in Table 4 . Stellite 21 can only be cast or applied

as a weld overlay and then ground. Haynes Stellite Alloy no. 6B is a wrought alloy. These two alloys appear to be representative of the best

of the stellite alloys in their respective categories for the com- bination corrosion/wear application being considered here. Heywood and

Womack( 2) suggest "Stellite facing" in their proposed valve design for

the high temperature air heater unit.

Resistance of Alloys to H

2 S

The update report on the Baseline Plant design by Jackson, et al ( 3),

at the recent 16th Symposium on the Engineering Aspects of Magnetohydrodynamics,

indicated that the proposed potassium seed recovery system will require

the handling of high- temperature gases containing high concentrations of

H

2 S. (See Figure 1 composed from data in the above paper by Jackson, et al.)

In respect to this requirement for alloys to resist attack by H

2

S at high temperature, some relevant data is given in Table 5a which was obtained in Phase I of the materials program of the Metal Properties

Council, Inc. for coal gasification plants. ( 4 ) Only their 1500° data,

which corresponds to the 1060°K reactor temperature (Fig. 1 ) are included in Table 5a . The alloys are arranged in the order of

increasing thickness loss. Although these absolute values of loss rates

cannot be assumed to apply in the atmospheres experienced in the MHD

application, the relative alloy ranking in respect to attack can

probably be assumed to hold. This assumption is supported by data

reported by Hungtington Alloys Division of INCO ( 5 ) on thickness loss

rate in a different ^S-containing atmosphere. (The actual data are reproduced in Tables 5b and 5c .) Although there are very few alloys

that are listed by both sources, the three that are, i.e. Incoloy 800

and type 309 and type 304 stainless steels fall in the same general

relationship. There probably is no basis to make a distinction between

types 309, 310, 446 stainless steels, 50 Ni-50 Cr and Incoloy 800 on the

basis of these data, but Incoloy 600 and 9 Cr-1 Mo alloy appear to be

distinctly inferior in their ability to resist H

2

S attack, with the

others falling somewhere in between the first 5 and these last 2 alloys.
.
71

References

1. ASME Boiler and Pressure Vessel Code, 1977 edition, Amer. Soc

Mech. Engrs., 345 E. 47th St., NY, NY 10017.

2. J. B. Heywood and G. J. Womack, Open Cycle M.H.D. Power Generation,

p. 38, Pergamon Press, Oxford, 1969.

3. W. D. Jackson, R. L. Lawit, R. A. Stoudt, M. G. Klett, J. C. Cutting,

and C. D. Maxwell, Status of the Reference Dual-Cycle MHD - Steam

Power Plant, 16th Symposium, Engineering Aspects of Magnetohydrodynamics, Pittsburgh, 1977.

4. A. 0. Schaefer, C. H. Samans , M. A. Howes, S. Bhattacharyya, E. R.

Bangs, V. L. Hill and F. C. Chang, A Program to Discover Materials

Suitable for Service under Hostile Conditions Obtaining in Equipment

for the Gasification of Coal and Other Solid Fuels, Annual Report

to ERDA 1/1/75 to 12/31/75, The Metal Properties Council, Inc.,

345 E. 47th St.

,

NY, NY.

5. "INC0L0Y Alloys", Huntington Alloy Products Division, INCO,

Huntington, WV.
'
Mechanical

Properties

of

Wrought

Alloys - n

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J= e 0 rH rH rH CO fa fa fa rr 3 o fa PH rH O H n rH CN CN rH rH rH 4J

4HrH

cj

r—.

U

U CJ CJ CJ CJ

u CJ

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cj cj o o o o o £ o o CO CO rH P P p P o o P p P P CN rH P P P 0 fa <P ro P p P» rr ^r fa fa *3* 3* 3* © P* 3 3 3

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(/)

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rH fa 0 rH rH U) 0 0 0 0 rH 0 0 0 0 4-> C/5 fa >4 fa rH fa fa fa fa fa 2 0 0 0 c c c C O c rH rH rH 0 E-* 0 0 0 2 < 2 H» 4J c 0 0 0 0 0 0 0 0 0 e fa c c rH P cn M 3 o fa fa c3 U5 U] >i o 0 0 o o u o o 0 •H < >4 >4 0 2 «d < o fa rH 3 fT3 (TJ rt3 c c c c c c c c c CD (0 (0 2 5H CJ

E-» CO CO CO 3 •H 2 2 2 H JH

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72 010

rHao

Table

2.

Chemical

Compositions

of

Selected

Wrought

Alloys

CO CO CO Z

ro co co

m OOOoa . . . O

CL di Qj m in in

a

<3* ^ o o o o o o o o

co in co

m

O in o ro o o o o o m o o o o o m

fN (N (N

H

(litO l l l l I l I I d d d fO X X XI X

(N (N ^ O H (N H m III. CO o o o ° o co r- M OJ H (N I I I I CO ^ id n * z z

Qj <N CN O O O cn d d > X X

in (N CN ro O O O O O

in * o O O 5T o O (N H i i d o ^ * m m • o d d o X x <n

in o

l£> COO

in * in • H • <3* I CN l CN

ro . O

O'

O IH H l I

CO BJ CO co o 'JP CO co o a <Jd o o JS o o H o o in f • o • • o o • 1 I o i * * 1 3 3 3 CO 1 in 3 i 3 3 1 u U U co 1 o u i u O 00 co o

in ro in ro ro N* o o o O o O O o o

inH

in O' CO co in in in in o 1 O o O o o H o ro H H in o o O o o o o O O O* in

in in * o o o o CN ro in in in d* 1 • o • 1 • 1 • • 1 H i o O o O o o CNo

in in in o in

in ro o CD CN CN o o H 1 O o O 1 H H

rH

O

d I X in

in o l CN H |c

in «h

in in E O O' O O i0 in m o* H O' o o CN H • * i t CO CN O' ^

CN CNHH

dHUH0 i

CO I I asox in O h • IM d CN X co ?oH0X

CD0

cn

10 o o

O' O'

CN CN

d

ro ro

X

10 10 H H I I o o

ro ro d d X X

H so c < 10 o 0 cn co cn •H H 0 CN \ X o > X 0 cn | o H H O o CN o 0 d 0 d CJ H r" co in r" r-' o o o in H c X CQ 10 o H H 10 10 CD CO CO D i—

cr CO >i >i in 10 10

X <

•H * • •H 0 0 in 73

cn cn CO CO CO CO z H

H

HHH rH

d

H >. >1 >1 0 3 cn • 1 H H cn 0 0 0 0 H 0 0 0 0 X cn Q 0 CO CO CO CO 0 0 0 0 G C g u c H H H 0 d H fc* X X c 0 0 0 0 0 0 0 0 0 £ co c rr o 10 10 cn cn >. 0 u u 0 0 u o o o •H U •H o H H 13 d d d c c c c G c c c c 73 X

O'

d ro ro ro o* E E H

H

HM

M

H

M

H

H

X

H

U-i

X •H

d 2 CO z X h cn ro o' CO O O H (N ro 'T in 10 o*

H O'

CN CN I I <T> O H CN

CN CO

CO CO

SC

H

cn cn 0 0 c c > >1 d d X X

CO ON

CO CO CO ro O O

CU 0* ro ro O O

in co H O O O O

i0 r- o o

73 • * CN o 0 0 in H CN l • - 1 1 t 1 1li li1 ( l ro ON ro O 1 11 ON co X • * * O ro o H d o d CD on in — ‘—

• • CN l l 1 11 1 o in • 0 1 1 t 1 H X 1 !( 1 l in t d

r* t*'*

M • • X <3* O cn

H

X • l 1 3 1 1 1 11 1 1 1 1 1 o cn 1 1 1 1 73c CN d in CO CO * cn in ro ro CO 11 ro t 1 i • • • 11 l t 1 1 1 O CN H O O O in i 1 t 1 0 i0 CN o

CO . in • H ON CN

ro ro

O 1 1 c CO CO 0 • • • cn rr •H 1 ro ro CO 1 1 1 1 O H H • • 'V i « 1 l 1 1 1 1 0 0 73 O o in I I I Iad S *3 ao ^ > s O cji Ho s n U O X
.
74

Table 3.

ASME Boiler and Pressure Vessel Code, Maximum allowable Tensile Stresses

Alloy ASME

Spec. No.

Shape

Max. allow, tensile stress, psi

@ 1200°F

(648°C)

@ Max. allowable temp,

(temp, in parenthesis)

304 SS SA 213 tube 6,000 1,400 (1500°F/815°C)

304L SS SA 213 tube 9,000 (800°F/427°C)

plate 11,000 (800°F/427°C)

310 SS SA 213 tube 6,000 700 (1500°F/815°C)

316 SS SA 213 tube 7,400 1,300 (1500°F/815°C)

446 SS SA 213 tube 14,300 (650°F/343°C)

Incoloy 800 SB 163 tube 7,900 1,700 (1500°F/815°C)

ff SB 408 bar 7,900 1,700 (1500°F/815°C)

Incoloy 800H SB 407 tube 8,400 2,500 (1500°F/815°C)

Inconel 617 (Has not been entered into the Boiler Code)

Hastelloy X SB 435 plate 11,300 4,800 (1,400°F/760°C)

SB 572 bar 11,300 4,800 (1,400°F/760°C)

Hastelloy C-276 SB 574 bar 21,800 (1,000°F/538°C)

Hastelloy G SB 582 plate 19,000 (1,000°F/538°C)

SB 581 bar 6,000 (1,000°F/538°C)

Haynes 25 (Has not been entered into t' ie Boiler Code)

Haynes 188 (Has not been entered into the Boiler Code)

9Cr-lMo SA 213 tube 1,500 1,500 (1,200°F/648°C)
V

.

'

_

.

.
Nominal

Composition

p<

LO

CNom3

rQ

CO

CN I CT\X3go

CNX3gO

CNX3go

cnmmI in

cnX3g

in3

JOO

cno

CM o cni

r-'- i CO

CN CN

4-V pH p—s CP 3 CN ^ cO w w 0) o 3 O H 13 H X 3 -H 3 -H 3 ,H >i 3 i—

3 iH O 3 iH O

>> 3 i-1 >1 3 pH 3 H>—

3 H rH X cn <3 X cn C

oo3

•HTO3

rJ

5-i 00 TO33

fn3

pH

5-i 0)>O

TOt—I , 0)&

>i

-50003

•H

CJ3

TO

5-i cd 30

5-i O

0) TOo

5-i H

CJ

0)

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5-1 01>oa

5-4 0)PuO

54 P4

cdo

-H3

cd jou

0) 2 oH

JO

003O

54 12

h cn o jo PQ Cd tH o S' 1 N g H H H4 CH 3o

•HH

cd003oI—

w

JOH

00

TO fj3 r-l ^ 3 5-1

H

•H

h CO >-rCO ^

a) 3 3rj 3 2 +

4J Cj ^ a) cn P HI

cnP cnP

3pJ

T3ho

cd

O

oo oi

a)

543H3

54

01a§

E—

o

tH

CN

CNvO

00 i-- O'! rH

40

CN CO o

oo cn CO -o 1—

CO

n o i—

<r i—

1—

i—iH

-Vo 00

on cn

CJ CJ / s CJ o CJ o o o C 1 O CO . O 43 o <r H in r— * CN kO CN '£> Pi rH ^ Pi i—l '—

i—

X ^ CN NO oN ' TO . 3 o 3 O

HIS

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3rH <u rH 3 HrH cd H i—1 00 cn< 00 cn <0

75 33

pH

cd>

cn

cn

a)c

TO

54

cd

a.

a)

CJx

a)l

5n

cd rH

543>O3

TO44O33

•H

543H3g

cn3

CJ3a

(d)

Data

on ^"

rolled

plate

material

*Workhardens

to

C45 +

Ksi=

1000

lbs/sq

in.;

to

convert

to

MegaPascals

(MPa),

multiply

by

6.895

76

Table. 5a. Corrosion Rates in H 2 S from The Metal Properties Council data (Ref. 4)

Alloy Loss of sound metal in./ yr .

310 S.S. .004

309 S.S. .006

50 Ni-50Cr .006

446 S.S. .008

Incoloy 800 .015

304 S.S. .039

316 S.S. .041

INCONEL 601 .657

INCONEL 600 4.9

*based on 1000 hr. test at 1500°F(815°C) in 1% H ? S, 24% H

2

, 18%C0 2

, 12%C0

,

5% CH , 1% NH q , bal. Ho 0, except Inconel 600 which was a 100 hr. test.

li J z

Table 5b.

Corrosion Rates in H

2 S from Huntington

Alloys data (Ref. 5) in 7.0% H

2

S

,

bal. H

2

.

Table 5c.

Corrosion Rates in H

2

S from Huntington Alloys data

(Ref. 5) in 1.5% H

2 S, bal. H2

Corrosion Rate, mpy

Alloy

Temperature,

• F 330-hr Test Long Term INCOLOY alloy 800 600 1.1 0.6 750 8.7 2.3 900 28.2 5.8 Type 321 Stainless 600 1.9 0.9 750 14.6 3.3 900 38.1 5.3 Type 304L Stainless 600 2.1 0.9 750 15.1 2.6 900 36.8 5.7

9 Cr-1 Mo 600 37.9 17.1 750 172.2 64.7 900 383.2 128.0

a 5 to 7% by volume hydrogen sulfide in hydrogen at 2000 psig.

** Calculated from slope ol weight-loss/time curve.

Alloy INCOLOY alloy 800

Type 309

INCOLOY alloy 804

Type 204 ELC

Type 202

Type 304

INCONEL alloy 600

Corrosion Rate, mpy

2.9

4.9

5.0

5.0

7.7 7.9

10.1

a Test duration, 730 hr Test temperature, 750' F. Hydrogen pressure, 485 psig HjS

concentration in H 2

. 1.5% by volume.
sy
Fig. 1 Schematic of Seed Regenerator Function (from Ref. 3)

Spent Seed

K

2

S0

4

'd

CD

(l) CQ

rC

cn

(DU

co o

o

CVJ

Fixed

Bed Reactor

1060°K —

0) > C •H

<H•H

CQ CQ

cd cd O bO

CQ i—

O •H cd CM

CA O K

<D O O

s o

a o <n >5 u CM

CQ W

h2s, co 2

Recyle Gas

1 ~

from Catacarb unit