O&S' iSy 7^» ' II a

E- (49-1) -3800-5

Distribution Category UC-90g

DEVELOPMENT, TESTING AND EVALUATION OF

MHD-MATERIALS

Quarterly Report

for the period July - September 1976

H. P. R. Frederik.se, T. Negas and S. J. Schneider

Inorganic Materials Division

Institute for Materials Research

National 3ureau of Standards

Washington, D. C. 20234

Date Published: September 30, 1976

PREPARED FOR THE UNITED STATES

ENERGY RESEARCH AND DEVELOPMENT ADMINISTRATION

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

DEVELOPMENT, TESTING AND EVALUATION OF

MHD-MATERIALS

H.

Quarterly Report

for the period July - September 1976

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

Inorganic Materials Division

Institute for Materials Research

National Bureau of Standards

Washington, D. C. 20234

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

State 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."
r

V v.i
Table of Contents PAGE

Abstract 1

Objectives and Scope of Work 2

Summary of Achievements: July - September 1976 (Completed Milestones) ... 3

Task G. Materials Program Coordination 4

Program Review and Consultation 4

US/USSR Cooperative Program 4

Closed Cycle Program Assessment 4

(Task H. U-02 Phase I - completed) 5

Task I. Operational Design Properties 6

Viscosity of Coal Slags 6

Electrical Conductivity (Slag and U-02 Materials 11

Vaporization (slag rejection and fractionation) 19

Task J. Corrosion and Diffusion 23

Phase Diagram of the System K 0-A1 0 -SiO 23

Diffusion of K into Al„C> 31

2 3

Task K. Materials Testing and Characterization 36

Spinel Electrodes for ANL-Reynolds Test 36

U-02 (Phase II)

Pre-test Characterization of Materials 40

U-25 Bypass-Channel (Plans) 43

Soviet Test at UTSI 43

Characterization of AVCO-test (August 1976) 43

Evaluation of Fluidyne Preheater Test Samples 53

Task L. Assessment of Materials for Downstream Components 63
4 i > i
1

ABSTRACT

During the 5th Quarter there has been strong emphasis on the

preparation of MHD-materials for upcoming tests, and characterization

of electrode and insulator materials that have been tested recently.

In the first category are the spinel materials for the U-02 test (Phase II)

,

and the electrode/insulator assemblies (Mg-Al-Fe spinels as conductors,

pure spinels as insulators) for the Reynolds test (under auspices of ANL)

,

both to be conducted during the latter part of September. Characterization

experiments (X-ray diffraction and SEM) are being performed on a variety

of electrodes that were tried out in the AVCO-Mark VI generator for

periods of 5 and 20 hours (with coal slag added to the fuel) . Also,

a new series of preheater samples, tested by Fluidyne, are being evaluated.

Calculations were made concerning the amount of various slag

components in the vapor phase ’under conditions corresponding to those of

a two-stage combustor and of a radiant boiler section (in the PERC-model

of an MHD- system)

Considerable progress was made in the determination of the phase

diagram of a simplified seed-slag system (K^O-Al^O^-SiO^ )

Important data were obtained concerning the diffusion of potassium

into alumina, and that of Fe from Mg-Al-Fe-spinel into pure Mg-Al-spinel

In connection with cooperative programs, N3S staff members visited

a number of other MHD contractors: Montana Energy Institute, Montana

Tech, and Montana State University, Westinghouse , AVCO, MIT, Fluidyne

and Reynolds Metals.

2

Objectives and Scope of Work

The overall objectives of this program are 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.

I 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 samples and structures that have been subjected to real or simulated MHD conditions.

The activities are grouped under six tasks:

G. Program Management Coordination (Assisting EKDA 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).

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 - July - September 1976

( Completed Milestones - see Work Statement)

During the 5th quarter of Fiscal 1976 the following program points

were accomplished in accordance with the projected milestones. In a

number of cases the actual achievements differ somewhat from the plans

as formulated a year ago, because of some necessary changes in work

assignments and schedules, made in discussions with, and agreed upon

by, staff members of the ERDA-MHD Office.

la. Planned and prepared negotiations of the US/USSR Working Group on

Testing of MHD Materials: Phase II and III.

2. Participated in reviews, planning sessions, preparation of working

documents in the area of MHD-materials development, (e.g., study on

Closed Cycle MHD)

.

4. Measured the viscosity ( and electrical conductivity) of a variety of

coal slags (in particular one Mid-Western and one Western coal)

.

20a. Determined temperatures of interfaces in one spinel-metalfeltmetal sandwich electrode under large temperature-gradient

produced by a CO^-laser.

29. Determined diffusion rates of alkali seed into various alumina samples.

39a. Assisted in the testing and pre- and post-test analysis of specimens

and electrode-assemblies exposed to simulated or real MHD conditions

in a variety of test rigs and generators. [U-02 (Phase II)

,

AVCO

(August test)

,

ANL- Reynolds]

.

Assessed the degree of hot corrosion attack on a number of steels

and alloy steels designated for use in the steam superheater and

reheater coils.

42a.

4

Task G. MATERIALS PROGRAM COORDINATION (S. J. Schneider)

Program Review and Consultation Activities

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

and in program review meetings at AVCO, Westinghouse and Reynolds Metals

(ANL U-25 Program) . The conclusions and recommendations resulting from

these meetings are reflected in Panel reports to ERDA or through direct

consultations with ERDA staff.

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

for review. During this period, NBS staff provided 23 such evaluations.

Arrangements and plans were also formulated relative to the various

testing and characterization activities carried out between NBS and

MHD-ERDA contractors (Battelle NW, Fluidyne , Westinghouse, ANL, UTSI

and AVCO) . These efforts relate both to the National Program and to the

US-USSR Cooperative Program. Special emphasis is placed on the transfer

of information and data to participating groups so that maximum impact

to individual development programs can be gained.

US-USSR Cooperative Program (Materials)

Major emphasis this period was focused on the coordination of the

various activities required to effect the Phase II U-02 in September 1976.

The test modules (anode and cathode walls) were delivered on schedule by

Westinghouse in August 1976. The test plan, worked out by Westinghouse

and NBS and approved by the USSR, forms the basis for a 100 hr. endurance

test under MHD (U-02 conditions) . The results from this test will provide

important information necessary for the selection of the electrode system

of the U-25 channel being developed through ANL management.

In preparation for the forthcoming October Working Group meeting in

Moscow, a tentative program for the U-25 high field loop has been formulated. Also, work continues on the materials section of the Joint US-USSR

MHD Status Report. In this context, the USSR input has been reviewed and

a suggested revision by USSR was requested.

Closed Cycle Program Assesment

In cooperation with RAND Corporation, NBS has been actively engaged

in the status and assessment of closed cycle MHD technology. S. J. Schneider

served as chairman of the June 28-July 2, 1975 ERDA-RAND Conference on

Closed Cycle MHD and additionally is providing input on a continuing basis

to the preparation of RAND assessment documents.

5

Task H. U-02 Materials Testing and Characterization (Phase I) (W.R. Hosier)

Final Report U-02 Test (Phase I )

The Soviet contributions to the report have now been incorporated.

The report will be submitted to the US/USSR Steering Committee for

approval in October.

Task H was completed as of June 30, 1976 .

6

Task I . OPERATIONAL DESIGN PROPERTIES

a. VISCOSITY OF COAL SLAGS (W. Capps and D. A. Kauffman, Inorganic Glass Section)

Introduction

In order to better coordinate the efforts of the many MHD contractors,

a decision was made to "standardize" on a typical midwestern and a typical

western coal. These are Illinois #6 coal and Montana Rosebud coals. This

will reduce the compositional range of coal slags to be encountered and

will facilitate comparisions of data, etc. Caution must be used, however,

because even these two coals cover a fairly wide compositional range and

two samples of the same coal can be quite different in composition.

Slag Preparation and Viscosity Measurement

Viscosities were determined for several real coal slags and of several

synthetic slags made to resemble slags whose analyses were available.

1) Melt No. K615 is a synthetic slag whose composition was made to

resemble the Illinois #6 slag analysis from the NASA ECAS specifications.

One deviation from this was the intentional omission of SO^. The viscositytemperature relationship is shown in Fig. 1.

2) Some flyash from the Bureau of Mines PERC combustor resulting

from combusion of Illinois #6 coal was melted. Viscosities are shown in

Fig. 1. This is K769.

3) A synthetic slag was made from an average composition from five

analyses of Illinois #6 coal slags listed in the Selvig and Gibson report.

K748 has no sulfur. Viscosity data is shown in Fig. 1 and is nearly

identical to data for K615 above.

4) A synthetic slag was made identical to K748 above with 5% SO^

added. This is K749 and is also shown in Fig. l. The SO lowers the

viscosity about 9% over the temperature range shown. Both K748 and

K749 have lower viscosities than the real ash (K769)

5) Some bottom ash from Montana Rosebud coal was obtained from a

commercial power plant. Several independent determinations of viscosity

were made and are shown in Fig. 2 as melts K603 and K624. Results are

quite similar.

6) A synthetic Rosebud slag was made from the NASA ECAS specifications

and contained a very high sulfur content. This is K612; the viscosity

is quite different from the bottom ash as shown in Fig. 2.

7

7) Another synthetic slag was made of the same proportions of the

oxides as in K612 with the SO^ left out in order to see what role the

SO^ might play. This is K620. Its viscosity is higher than K612

showing that here too the SO^ lowers the viscosity. See Fig. 2.

However, even without the SO the viscosity is much lower than the

real bottom ash (K603 or K624)

S) Flvash from the same Montana power plant as K603 and K624

formed by burning Rosebud coal was measured. This proved very difficult

as the melt out-gassed for a long period at 1400 C causing a foamy crust

to form which was too viscous to allow the bubbles to escape. At 1600 C

the crust was fluid enough to permit the foam to collapse and allow the

melt to clear of bubbles. The viscosity is shown in Fig. 2 and is similar

to the bottom ash from the same coal.

PLANS

1. Measure viscosity of a MHD channel slag from the University of

Tennessee Space Institute from Illinois #6 coal.

2. Make a synthetic slag having the composition of the Russian

Krushnitz coal slag. Measure viscosity.

3. Try to identify cause of excessive bubbling of Montana Rosebud

power plant flyash.

4. Continue study of role of SO^ in determining viscosity.

Analysis

furnished

by

Montana

State

University. 8

Table

1.

Coal

Slag

Composition
i'
9

o

o

c\ in CO <T> VO r-i p* VD

*4

/

1 1 lx Q

(sas-pod

°"t'5o-[) Aq.TsoosxA

o

o

CN

Figure

i.

Viscosity

of

Illinois

#6

Coal

Slags

K769

Combustor

Flyash,

Pittsburgh

Energy

Research

Center,

Bureau

of

Mines.

K615

NASA

ECAS

Specifications.

K748

Average

Composition

from

Selvig

&

Gibson

of

5

111.

#6

coals

(No

S03

)

K749

Average

Composition

from

Selvig

&

Gibson

of

5

111.

#6

coals

(5%

SO^)
K

‘

i
10

Viscosity (log poises)

NJ

toOO

H- z z o o

NJ * * * * * cQ

05 <n -J 05 05

C

H* NJ NJ fO

o S< NJ o CD u> CD 11 NJ 11j Z z S3 S3 i > > 0 0 < M ! CO CO < 3 3 H- u> :

> > QJ 3 3 CO O j CO 0 O O ° i H M 3 3 3 0 n n CO 1 > > Hi > > H* » c

co

S C0_ CO 3 0 ? 3 i CO co 3 j S3 S3 Hi Hi O

CD 0) 2 H If Hi ! o 0 0 O 0 i H- H- 3 3 3 2

Ml Hi rrO

H* H- 3 2 2 3 j O O 3 0 0 3

0) 0) 3 3 3 0) 3 rr 3 rr 3

H* H- >3 3 o> 0/ I 0 0 0 3 3 M ' 3 3 s: 0> 03 S3

CO CO (D 0 o

rf S3 S3 CO

O ; ,—, 0 O CD \

s: s: 3 s:

S 3 i-3 * H- H- l—1 (D CD C CD - 3 3 0) if If Cu 3 3 3“ 3 S3

rr S3 S3

n CD NJ 3 • M M 0

'J 0 0> 3 0) 0J 1 • 3 3' M 3 ! U C/5 3 rr C t dP o CO if i u> I-1 CD cn • „— 0) O co co ua .—

U) OJ 3 CO

Q

—

3 3 • —

• S3 S3 3 M h-> U1

—

CD CD

o 3 E/ > O ooo A£

N>

S'

/NJ o

U)1 oo
tt
11

b. Electrical Conductivity (A.J. Armstrong, W.R. Hosier, H . P . R , Frederikse)

During this reporting period the electrical conductivities of two representative coal slags were investigated. Conductivity measurements were also

performed on 5 electrode materials that will be tested in the U-02 MHD facility

in September. In addition, extensive changes were made in the cooling system

on the ASTRO high temperature-controlled atmosphere furnace.

Coal Slag . The two coal slags whose electrical conductivities have been

measured in air are Montana Rosebud and Illinois # 6 . (Figs. 3 and 4). The

Montana Rosebud slag is from power plant bottom ash obtained via Mr, W. Capps

of NBS. The Illinois #6 slag was removed from the generator walls of the UTSI

MHD facility and was obtained from Dr. L. Crawford of UTSI. The conductivity

of the Rosebud slag below 1150 °C is much lower than that of other coal slags;

in that range the conductivity falls on the very edge of the "allowable window"

as calculated by R. J. Rosa.-*-

U-02 Materials . The electrical conductivities of the two NBS spinel electrode

materials (both were 3 MgAl 203 : ^

6304 ) are shown in Figs. 5 and 6 . The results

agree closely with those measured previously on this same material and reported

in our previous quarterly report. It should be realized that the magnitude of

the conductivity depends somewhat on the structural properties (porosity, grain

size, etc.) determined by the particular fabrication method.

Of the three other U-02 electrode materials investigated, two were LaCrO^

with approximately 5% MgO added. (Westinghouse and General Refractories)

.

Both samples evidenced somewhat higher electrical conductivity (Figs. 7 and 8 ) than that reported for pure LaCr03 by Grabner, Hosier and Frederikse. 2 The

third electrode material, 75% hercynite + 25% magnetite, (developed by MIT)

showed a marked change in electrical conductivity in air below 1300 °C (Fig. 9)

.

As this was not seen in 10

-

3 and 10-

6 oxygen partial pressures obviously

oxidation and reduction has a lot to do with this abrupt change.

Watercooling for the ASTRO Furnace

During the past half year problems have arisen with the cooling coils of

the ASTRO- furnace . 3oiler scale has built-up on the hot surfaces in contact

with the hard cooling water. After considering various deflocculants or filters

it was decided to replace the water-cooling system with a closed one using

distilled water and a heat exchanger. In this heat exchanger the continuously

circulating supply of distilled water is cooled by chilled water of the central

laboratory system.

Future Work : 1. The electrical conductivity of spinels with various amounts of Fe.,0^ will

be investigated.

2. The conductivities of other eastern and western coal slags will be measured.

3. The effect of specific admixtures on the conductivity of coal slag will be

explored.

^R. J. Rosa, AVCO Monthly Report of Office of Coal Research, January 1975.

2

L. H. Grabner, W. R. Hosier and H. P. R. Frederikse, Proc. of the 14th Symp.

on Engineering Aspects of MHD Power Generation, Tullahoma, TN, April 1975.
:
CONDUCT

IVITY(ohm-cm

)

1

12

Fig.

T(K)

3. Electrical conductivity of coal slag (Montana Rosebud) as a function of

temperature measured in air.

CONDUCTIVITY

(ohm-cm)

1

T(K)

^

Electrical conductivity of coal slag (Illinois #6) as a function of

temperature measured in air.
,
14

Fig.

5.

Electrical

conductivity

of

porous

spinel

75

MgAl^

-

25

Fe^

as

a

function

of

temperature

at

several

partial

pressures

of

oxygen.

(BNW)

.

Figure

.6.

Electrical

conductivity

of

sintered

spinel

e

and

various

oxygen

partial

pressures.

|?f,5

1060

903

7RO

679

597

16

T(K)

Fig.

7.

Electrical

conductivity

of

LaCrO,:Mg

(G.R.)

asa

function

of

temperature

as

pressure.

17

2<^>ON(X) lr) m- to cm

, (UJO-UJ'JO) AllAIiOnaNOO

T(°K)

Fig.

8.

Electrical

conductivity

of

LaCrO^:Mg

(Westi

nghouse)

as

a

function

of

temperature

and

pressure.

IU

I

,

iwyj

18

(ujo.iuqo) AllAllOnaNOD

19

c. Vaporization Studies (E. R. Plante)

During this quarter it was requested that NBS analyze the problems

of combustor slag rejection, slag f ractionation and the general problems

of slag composition as a function of wall position in an MHD channel and

downstream components.

At the present time, a difference of opinion exists concerning the

optimum value of ash carry over. Some consider that the liquified coal

ash components can serve as a continually renewable protective layer for

the electrode structure of an MHD channel which will protect the electrode

structure from erosion, corrosion and destructive arcs while not appreciably degrading channel performance. As a matter of fact, coal slag is

considered as a potential channel replenishment material. If true,

insufficient slag carryover could lead to bare spots and degradation or

electrode structures.

Others however, believe that generator performance will of necessity

be impaired by the presence of liquid ash because of short circuiting of

electrode structures. In addition, the problem of seed recovery is

expected to be more difficult if the seed has the opportunity to react

with large amounts of coal ash. Finally, the amount of dust and fly ash

in a system with large percentages of ash carryover are seen as difficulties leading to possible plugging problems for air preheaters.

In view of these problems, it has generally been considered desirable

if possible to design and operate a combustor which is capable of rejecting

90% of the ash.

A review of MHD combustor designs was presented by Henry, Kurtzrock

and Bienstock (1) who advocate the idea that high ash rejection is desirable.

As a limitation on ash rejection in a one stage combustor, they state that

"Equilibrium calculations (with the technique of Spencer, 1973) of combustion of coal (a Pittsburgh seam bituminous coal) at a pressure of eight

atmospheres with about 90% of stoichiometric air preheated to a temperature

of 1280K with 6% heat loss predict a plasma temperature of about 2660K.

At these conditions, more than 25% of the coal ash is predicted to be

vaporized. This puts an apparent upper limit on slag rejection capabilities

of single-stage MHD combustors of 75%.

The purpose of this discussion is to outline some of the concepts and

thermodynamic reasoning that can be called into play in such design considerations .

The actual combustion process is extremely complicated. The rate of

vaporization of slag particles may be limited by the rate of diffusion of

vapor away from the particle, or by the rate at which heat flows to the

particle. In addition, it will depend on the particle composition.

20

At any specified temperature, pressure, and oxygen/fuel ratio one can

estimate the fraction of slag that would be in equilibrium with a condensed

slag phase if some assumptions concerning the activity of slag components in

the condensed phase are made. Variations in the basic assumptions are

possible. For example, the ash components in the coal feed stock are

probably not homogeneous so that more ash components might be present in the

vapor phase than would be if establishment of equilibrium with respect to a

hypothetical ideal solution were achieved. In addition, ash components

can be transported to the vapor phase via formation of a number of different

types of reactions and equilibrium with respect to each reaction may not

necessarily be established.

The metals from principal ash components can be carried in the vapor

phase as a variety of species as indicated in Table 1.

These species were assumed to be formed in the calculations.

Table 1

Ash Component Vapor Species

Si0 ? SiO, SiO^, Si

MgO Mg, MgO, MgOH, Mg (OH)

2

CaO CaO, Ca, CaOh, Ca(0H) ? FeO Fe, FeO, Fe(0H)

9

A1

9 0

3

Al, A10, A10 , A100H, Al 0, A1 (OH) , A10H , Al 0

2

For purposes of this report and following the method adopted by Spencer

and Orning (2)

,

we assume that the slag components are present in an ideal

liquid solution and that equilibrium is established with respect to all

reactions. Having calculated the amount of slag in the vapor phase it is

assumed that most of the slag remaining in the liquid phase is effectively

removed to the cooled sidewalls by the swirl action of the combustor.

Thermodynamic data used in our calculation are from JANAF (3)

,

using

tables furnished through December 31, 1975 except for data for A1(0H)

? which was taken from Jensen and Jones (4). The data for AKOH)? was

obtained by flame photometric spectroscopy and has not been confirmed by

other laboratories. For example, Farber et al (5) were unable to detect

the A1(0H)2 molecule by mass spectrometric observations of Al in H2 /O

2

flames where it presumably would be a major Al-bearing species. Rejection

of A1(0H)2 as a possible Al-bearing species would decrease the volatility

of AJ^O^ signi-f icantly . It would not however have a significant effect on

the conclusions drawn from this calculation.

21

The variation in the amount of vaporization of each slag component

depends on the relative importance of each of the vapor species. The

greatest change in the amount of vaporization as the 0/F ratio changes

occurs for SiO n , followed by FeO and MgO while the effects are relatively

small for CaO and AI 9 O

3

. This is due to the relatively more important

transport of slag components by SiO, Fe, and Mg gases while a significant

amount of transport of CaO and AI2 O

3 is due to formation of hydroxide species.

Whether or not there is significant fractionation of the coal ash

will depend to some extent on the temperature at which slag separation in

the first stage of the combustor takes place and the fraction of slag

carried as particulate matter. The amount of Si02 ? FeO, and MgO transported to the vapor phase depends much more strongly on temperature than

does the transport of CaO and AI 2 O

3

. Thus, if the separation takes place

at a relatively low temperature, such as 2300K, fractionation effects will

not be very great. Any f ractionation that takes place because of the

vaporization processes will tend to be reduced because of carryover of

particulate material.

The results obtained here generally confirm the concept that large slag

rejection can be achieved in a two stage combustor provided that a high

percentage of particulate matter can be removed by the swirl motion of the

combustor. Some typical values for ash rejection based on a Montana Colstrip

coal are shown in Table 2. For this calculation the Montana Colstrip coal

composition is approximated on a molar basis as: C=1.0, H2 =. 393 , N2=.0046,

H2 0=.0207, 0

2 =.0815, Si0 2 = . 0195815 , Al20 3 = . 00582960 , CaO . 00590433 , FeO

.00388640 and MgO= .00272795 . This calculation assumed that preheated air

having a nominal N2 /O 2 molar ratio of four of sufficient quantity to give

an overall 0/F ratio of 0.6 or 1.0 was burned at a total pressure of 10 atm

or 1 atm to achieve a final temperature of 2300K.

Table 2

Amount of Slag in Vapor Phase at Equilibrium at 2300K

Percent Vaporized

Component P=10 , 0/F= .

6

P=1 ,0/F=1.0

Si0

2

4.3 1.8

FeO 4.0 4.1

Al

2 °3

2.0 .66

CaO 1.7 2.8

MgO 2.8 2.1

Vaporized 3.3 (wgt) 1.9 (wgt)
'
22

The conditions P=10 atm, 0/F=0.6, T=2300K are approximately what one

would expect at the exit of the first stage of a two stage combustor while

the conditions for P=latm, 0/F=1.0, T=2300K are approximately what would

be found in the diffuser-radiant boiler section of an MHD system. Since

the walls of the system will in general be below 2300K, nucleation and

growth of liquid slag particulates in the boundary layer region of the

channel can take place and slag will be transported to the wall surfaces.

Because of the relatively long life of slag on the channel walls as

compared to slag particulate matter in the core flow, the possibility that

significant variation in composition of channel slag can take place seems

remote. Rather, channel slag is likely to approach a "steady state" composition which will shift slowly in composition as variations in the input

coal ash composition take place.

Future Work :

1. Vaporization measurements on selected K^O-containing slag will be

resumed

.

2. Additional calculations of slag vaporization problems will be carried out

if necessary.

References :

1. J. M. Henry, R. C. Kurtzrock and D. Bienstock, Proc. NSF-OCR Engineering

Workshop on MHD Materials, MIT, Nov. (1974)

.

2. F. E. Spencer, Jr., and A. A. Orning, U.S.B.M. R.I 7786 (1973).

3. JANAF Thermo chemical Tables, Dow Chemical Co., Midland, Mich.

4. D. E. Jensen and G. A. Jones, J.C.S. Faraday J58, 259 (1968).

5. M. Farber, R. D. Srivastava, M. A. Frisch and S. P. Harris,

Faraday Symposium 8, High Temperature Studies in Chemistry, London, (1973)

.

23

Task J. Corrosion and Diffusion

a. Seed-Slag Interaction (L. P. Cook, C. L. McDaniel, T. Negas)

Phase Equilibria in the System K O-Fe^O^-Al 0 -SiC^

The Join KAl02 -KAlSiC>4 . Experiments using KA10 2 -Si02 mixtures continued.

The rapidity with which phase equilibria data can be gathered is limited by

the sluggishness of subsolidus reactions; for example, compositions heated

at 1400 °C must be reground and reheated four or five times before no

further changes in the x-ray powder patterns occur. A similar number of

heat treatments appears to be necessary at 1600 °C. Surface energy and

strain energy may provide most of the driving force for reaction, since

compositional differences of reacting phases are small in the vicinity of

(KA10 2

) c;q (SiO ? ) 5 q. Metastable equilibria also appear to be important.

Phenomena such as metastable melting and metastable persistence of low

temperature solid solutions upon heating have been observed.

Crystalline phases occurring along the join KAlC^-KAlSiO^ are:

hexagonal KAlSiC>4 (occurs in several polymorphs; reported by other

investigators at low temperatures but not encountered in this study)

;

"low temperature" orthorhombic KAlSi04 , "high-temperature " orthorhombic

KAl3iC>4, tetragonal Kn-xA^l+xS;''l-x04 ' cuki- c Kl-xA-*-l-xS:'-x 0 2 > non-cubic

KA10 2

. Crystallographic properties of these phases were summarized in

the previous quarterly report (Ref. 1) . Tentative limits of solid solution

at 1400 °C and 1600 °C can be set on the basis of experimental data

accumulated to date. At 1400 °C, "low- temperature" orrhorhombic KAlSiO^

ranges from ^48.5 to ^51.5 mole % KA102

, tetragonal Kj_ +xAl]_ +xSi ]__ x04

ranges v-56.5 to v58.5 mole % KA102

, and cubic Kj__ x

Al]__ x

Si x02 solid

solution extends to ^20 mole % Si02

. At 1600 °C, "high-temperature"

orthorhombic KAlSiC^ ranges from %47.5 to v51.5 mole % KA10

2

, tetragonal

K]_ +xAli +xSi-i _ x04 ranges from v55 to v56.5 mole % KA102

, and cubic

K. Al Si C> solid solutions extend to v25 mole % SiC> . Experiments

1-x 1-x x 2 2

thus far have not succeeded in dissolving more than 25 mole % Si0 2 in

the cubic solid solution. The extent of solid solutions at temperatures

below 1400 °C are difficult to measure, and it may be necessary to resort

to hydrothermal equilibration if phase equilibria data at temperatures

below 1100 °C are needed. Metastable " low- temperature " orthorhombic

KAlSiOd solid solutions containing up to ^56 mole % KA102 exist briefly

at 1400 °C before decomposing, and it is probable that the solid solution

field of this phase widens in the direction of KA102 at temperatures

below 1400 °C.

The nature of the phase transition occurring in orthorhombic KAlSiO^

at 1400-1450 °C is uncertain, hence, the designation "high-temperature"

and "low-temperature". Data accumulated thus far suggest it is reversible,

but that a compositional effect may be involved. If it is reversible, the
.
24

fact that increasing temperature favors a phase with doubled unit cell

requires an explanation - conclusive determination of the phase diagram

in the vicinity of KAlSiO^ will require a thorough understanding of the

structure of the phases involved.

Liquidus temperatures in the vicinity of KAlSiO^ have proved to be

higher than suggested by previous work (Ref. 2) . Our experiments to

date suggest a liquidus temperature at ^1850 °C for composition

(KAIO 2 ) 5 q

(Si 02 ) 5 q. The temperature at which melting begins for compositions intermediate between cubic and tetragonal solid solutions appears

to be in the range 1850 - 1870 °C. With the experimental set-up employed

(sealed Pt^gRl^g capsules heated inductively in an iridium crucible)

,

quenching was difficult, and a combination of data from x-ray diffraction,

light microscopy and SEM examination has been employed. The latter

technique (Figs. 10a, b and 11a, b) has proved to be especially useful in

this regard. Whereas KAlSiO^ melts have been cooled to a glass, melts

richer in KA10 0 appear to crystallize more readily.

Experiments in the Ternary System K^O-A^C^-SiC^ • Results of TGA and

mass spectrometric vaporization studies on crystalline KAlSiO^ were given in

a previous quarterly report. (Ref. 3). Two unexplained features were noted:

1) The failure of the three phase mixture KAlSi04-" 8 -Alumina"-KAlSi 20g

to

buffer Pj, and, 2) the fact that the vacuum decomposition reaction of KAlSiO^

involves KAlS^Og and " 8 -Alumina", rather than KAlS^Og and Al^C^, as the

equilibrium diagram (Ref. 2) suggests should be the case. 1) could in part

be explained by the presence of local concentration gradients in phases of

variable composition (both KAlSiO^ and " 8 -Alumina") . To check the possibility that 2 ) indicated an error in the published equilibrium diagram,

subsolidus experiments were conducted on mixtures of 33 mole % KAlSiO^/67

mole % CX-AI 2O 3 at 1500 °C for approximately two weeks, with periodic grinding. Results showed formation of a small amount of KAlS^Og and " 8 -AI 2 O 3

" initially, which disappeared with subsequent heat treatment. Further

changes in the x-ray powder patterns with continued heat treatments showed

only annealing of KAlSiO^ and C1 -AI 2 O 3

. These results are interpreted as

verification of a tie line between KAlSi04 and AI2 O3 at the expense of a

tie line between " 0 -AI 2 O 3

" and KAlS^Og. The decomposition of KAlSiO^

reported above may represent a metastable phenomenon. The topological

relationship of this decomposition reaction to the more stable decomposition

reaction, which could be prevented by the kinetics of AI 2 O 3 nucleation and

growth, is portrayed schematically in Figure 12.

The Join KFeC^-SiC^. A study of the subsolidus relationships on the

KFe02 -SiC>2 join has been initiated. A series of nine KFeC^/SiC^ mixtures

have been prepared under dry-environment conditions. While KFeC>2 is

extremely hygroscopic, reacted mixture of KFe0 2 /Si02 are not. Preliminary

experiments indicate at least four subsolidus phases exist in the system.

The low temperature form of KFeSiC>4 was indexed with a hexagonal cell,

a = 5.28/1 and c = 8 . 82A. This compound has a phase transition at 945 °C,

forming a crystalline phase which at the present has undetermined symmetry.

Complete melting was observed at 1100 °C for compositions rich in Si0 0

and 1400 °C in the silica-poor region. Future experiments will concentrate

on the crystal chemistry and the liquidus temperatures of the crystalline

chases

.
I

•:
25

Container Materials for High Temperature Liquidus Studies . Special

techniques for liquidus studies along the join KAlO n -KAlSi0

4 are clearly

necessary in view of the facts that temperatures above 1900 °C are required,

noble metal capsules are not capable of reaching these limits, and encapsulation is essential to prevent "K 20 " loss. Use of sealed refractory metal

tubes (e.g., Ta) in an evacuated high temperature graphite furnace seems to

be the best possibility. Calculation of one possible capsule corrosion

mechanism at 2000 °C has been carried out:

5 SiO, (i) + 2 Ta (s ) -* 5 SiO(g) + Ta 0

5

(a)

Corrosion according to this mechanism should not be serious. However,

reaction with KAlC^-containing mixtures is more difficult to estimate,

and an experimental evaluation will be necessary.

Capsule reaction with iron in the KFeC>2~KFeSi04 system is severe, but

it is hoped that use of Ag^Pd^ capsules will eliminate this problem at

elevated temperatures.

Implications for MHD Materials Research . Potassium aluminosilicates

are highly refractory materials, and in the presence of a K-rich vapor phase

can be expected to be stable in the presence of silicate slags up to

temperatures approaching their melting points. The implication of this is

that slag accumulating in the MHD channel under operating conditions may

contain crystalline KAlSiO^ and related phases. For iron-rich slags, the

chances of their occurrence is diminished, however. The effect of a crystal/

liquid mixture on such properties as electrical conductivity and viscosity

may have to be examined.

Attention should be called to the possibility of using KA1C>2-Si02

phase equilibria to deduce temperature gradients during MHD runs. Reactions

in this system (especially reactions of the homogeneous type) are sluggish

enough to prevent back-reaction, provided shut-down is not exceptionally

slow. Once the equilibrium diagram has been determined, possibilities should

be apparent. Likely candidates are the hexagonal ^ orthorhombic transition

and the "low- temperature" orthorhombic -»• "high- temperature " orthorhombic

transition.

Determination of phase relations in the systems KAl02~KAlSi04 and

KFe02~KFeSi04 should provide a starting point for the systematic extension

of seed/slag interaction studies to cover other systems. In particular,

it is suggested that the relative effects of increasing mole % CaO , MgO

and FeO on KAlSiO^-rich compositions can be seen by looking at the systems

Ca0-KAl0

2

-SiC>

2

, Mg0-KAlO

2

-SiC>

2

and Fe0-KA10

2

-SiC>

2

, respectively. Preliminary

studies in the system Ca0-KA102~Si02 (Ref. 4 ) have shown that it is ternary

and that increase in mole fraction CaO causes a significant increase in a^

for compositions with a slight excess of AI2O3. This seems to have 2

been corroborated by the few available data on K2O vaporization from high

CaO slag (see Ref. 4 ) and by available information on K^O absorption by

high CaO slag under operating MHD conditions (Ref. 5 ) . The systems MgOKA102~Si02 and Fe0-KA102- Si09 are ternary over much of their extent,

according to existing data (Ref. 6) . However, calculation of the high

temperature tie line distribution in the more MgO and FeO rich parts of

26

the systems suggests a marked difference from the CaO-KAlC^-SiC^ system

in that tie lines do not extend from the olivine phases to cubic KAIO2

solid solutions. The effect of increasing X„ _ and X ^ on aT, ^ may not

FeO MgO K 0

7

be as great, therefore. Figure 13 shows data presently available for

these systems.

References

1. H. P. R. Frederikse, T. Negas and S. J. Schneider, 1976, Development,

Testing and Evaluation of MHD-Materials : NBS Quarterly Report

January-March 1976.

2. E. M. Levin, C. R. Robbins and H. F. McMurdie , 1964, Phase Diagrams

for Ceramists: Columbus, OH, The American Ceramic Society, p. 156.

3. H. P. R. Frederikse, T. Negas and S. J. Schneider, 1975, Development,

Testing and Evaluation of MHD-Materials: NBS Quarterly Report,

October - December 1975.

4. L. Cook, E. Plante, T. Negas and R. Roth, 1976, Crystallization and

Vaporization Studies on Synthetic Coal Slag Compositions: Proc. 15th

Symp. Engineering Aspects of MHD , Philadelphia, PA.

5. J. K. Koester, R. H. Eustis, and M. E. Rodgers, 1976, In-Channel

Observations on Coal Slag: Proc. 15th Symp. Engineering Aspects of

MHD, Philadelphia, PA.

E. M. Levin, C. R. Robbins and H. F. McMurdie, 1964, Phase Diagrams

for Ceramists : Columbus, OH, The American Ceramic Society, p. 270.

6 .
I
27

Figure 10a. SEM micrograph (1000X, 20 KV) of (KA10

2

) (SiO ^

)

3 q

held at 1880 °C

for 10 minutes following pre-equilibration at 1600 °C. Recessed

area with deudritic structure is interpreted as devitrified melt.

Figure 10b. Enlargement (20,000X, 20KV) of area in upper left of Figure 10a.

showing triple grain boundary junction with recessed median

suggesting intergranular melt.

Figure 11a. SEM micrograph (1000X, 20 KV) of (KA1C>

2

) (SiC> 2) 4 q

held at 1850 °C

for ^1 minute following pre-equilibration at 1600 °C. Note

rounded grains (mostly tetragonal phase) , apparently produced by

local melting and recrystallization in response to surface energy

minimization requirements.

Figure lib. Enlargement (20,000X, 20KV) of grain boundary in Fig. 11a.

showing no evidence for intergranular melt, in contrast to

Fig. 10b.

Figure 12 .

10

Schematic y -10 /T diagram showing topology of some univariant K2°

potassia vaporization reactions of interest. Reactions are

uni variant only for fixed compositions of phases. Potassia

chemical potentials corresponding to the invariant point and

above apparently are not attainable under low-pressure experimental conditions.

Intersection of dotted line (Constant T) with reaction lines is

indicated by points labelled y-^ and Failing nucleation

and growth of a-Al^O^, represented by y , vaporization decomposition of KAlSiO^ would be controlled by the metastable extension of reaction involving its decomposition to KAlSi 0 +

" B-Al^O ^ "

r

represented by y^-

30

M«0

FeO

Fig. 13 . Phase relations known to date in the

systems MgO-KAlO^-SiO^ , FeO-KAlO^-

3iC>

2

and Ca0-KA1C>

2

-Si0

2

(Refs. 4, 6).

MgO, FeO-KAlSiO^ tie lines calculated

to be stable at expense of

Mg SiO , Fe SiO^-KAlO^ tie lines.

CaO

31

b. Diffusion

a) Diffusion of K into Al 0 (E. N. Farabaugh and J. R. Manning)

Three more diffusion couples were made during this period. All

were produced at 1400 °C for 240 hrs . They were ISOA-225 AI 2 O 3 (^83%

dense), ISOA-242 AI 2 O3 (^90% dense) and single crystal AI 2 O 3, each was

run in its own pellet of B-AI 9 O 3

. Only the single crystal material has

been analyzed to date. Even with the 240 hr. diffusion time, K was

detected only to a depth of 2-4ym from the surface of the single crystal.

The data was acquired in 2ym steps and after two readings, the level of

K X-rays dropped into the background. This time of 240 hours should be

sufficient to generate reasonable profiles in the other two Al^O^

ceramic specimens.

b) Examination of Diffusion in Magnesia-Aluminum- Iron Spinel Electrode

and Magnesia-Aluminum Spinel Insulator Material (E. N. Farabaugh,

J. R. Manning and A. J. Armstrong)

3ecause of the current interest in MgAl 20 ^-Fe 30^ solid solutions as

electrode materials, measurements of interdiffusion between MgAl 2 C>4 insulator material and MgAl 904 -Fe 3 C>4 electrode material were made.

Sandwiches of electrode-insulator-electrode were prepared, heated and examined to track the diffusion of Fe from the electrode into the insulator.

Since the iron concentration controls the electrical conductivity in

these materials, diffusion of iron from electrode to insulator could

strongly affect the operation of both the electrode and insulator in

operating MHD channels.

Fig. 14a shows, at 50 X, the insulator (left side) electrode (right

side) interface. This sample was produced after diffusion at 1600 °C for

4 hours in the absence of an electric field. Note that the iron spinel

is more porous than the pure spinel. Thus, there was no problem in

defining roughly where the boundary was. Fig. 14b is an Fe map of the

same area showing slight penetration of Fe into the insulator.

A detailed Fe concentration profile was obtained by collecting data

at intervals of 150ym, in the region far from the interface, to 20ym in

the regions of the interface to trace the diffusion of Fe into the

insulator. Fig. 15 shows the plotted data. The concentration of Fe is

constant in the electrode until the interface is approached. Then, the

concentration drops then rises again just before we enter the MgA^C^

insulator. The electrode-insulator interface is indicated by the arrows.

The Fe has penetrated ^0 . 3mm into the insulator while the electrode

shows a depletion band of M3. 4mm as measured from the curve.

32

Fig. 16 shows the data collected from a 51 hr test at 1600 °C with no

electric field. In this specimen we were able to traverse both electrodes

and the insulator. Again, we see the build up of iron in the electrode

just outside of this electrode-insulator interface. Here the penetration

of Fe is ^0 . 7mm from each of the insulator-electrode interfaces. The

depletion band in the electrode is also ^0.7mm.

Because of the observed difference in porosity, the transition from

electrode to insulator can be seen on the SEM viewing screen while the

data is being collected, so the position of the apparent boundary can be

easily and accurately established. We say apparent because there is the

possibility that some grain growth or sintering has occurred between the

two spinel pieces. This would make the apparent boundary move into the

insulator if the electrode material grew at expense of the insulator

material. It would be possible that the peak in the Fe profile appears

where the boundary originally was and was produced by Fe moving to this

boundary in a manner comparable to impurities moving to the boundaries

of grain in grain growth processes. Alternately, there may be a significant surface barrier preventing or retarding interdiffusion at the

boundary. One way to help examine this apparent boundary-related effect

would be to do experiments with the original boundary marked with fine

wires. The fact remains, however, that there has been an Fe pile up in

the boundary region. This effect can only enhance the lifetime of the

insulator material in an actual test run.

Further investigations of these spinels are planned including

testing the insulator-electrode 'sandwich' at different temperatures and

with an applied electrical field.

Calculations of diffusion kinetics in MHD electrode materials have

been oriented to concentrate on diffusion in these spinel materials.

Diffusion of probability estimates for interdiffusion of Fe and A1 on

the octohedral spinel sites are being made.
'
33

Fig- 14a. 50X secondary electron

micrograph showing the

MgAl 0. insulator on left

2 4

and Lron MqAl 0 „ on the

. . . y2 4 right.

toiol

34

3000*-

I I

|

•Uctrod*

i i ' i1 i

intulo<o<

1 l i . i ) I I 1. 1 A X-i—

^igure lb. Iron profile taken for 4 hr. test.
I
lolul

count*

35

' •

\ 4 . ...

m

. •

'

* , '• •

‘l /

/

/

1

•

i •

j_ V / — v — initiator *—

• / i \ / •

- 1mm \ y

* _

_| .

L. It . * > l 1 . 1 1-i- i ... i

Figure 16. Iron profile taken from 51 hr. test.
'
36

Task K. MATERIALS TESTING AND CHARACTERIZATION

a. Spinel Electrodes for a 100-hour Proof Test (W. R. Hosier, H.P.R. Frederikse)

In the early months of 1976, NBS was asked to supply a number of

Mg-Al-Fe-spinel samples for testing in the Reynolds MHD test rig under

auspices of Argonne National Laboratory. This material is one of the

spinels (described in an NBS paper presented at the 15th Annual Symposium

on Engineering Aspects of MHD, May 1976, Philadelphia) that holds promise

as an electrode material, because of its reported corrosion resistance to

slag and seed, reasonably high melting point, and good electrical conductivity between 500 and 1900 °C. Originally, the required electrode

assembly, [including preparation of the electrode ceramic, mounting in a

suitable electrode holder (with watercooling and electrical lead-outs)

and assembling with the appropriate interelectrode insulators] would be

designed and built by one or another of the several ERDA-contractors

involved in MHD development. However by early May it became obvious

that NBS itself had to assume responsibility for the construction of the

electrode-insulator test assemblies.

It should be realized, however, that NBS does not have the industrial

fabrication facilities required for building such a device. Although this

electrode/insulator system is relatively small, one needs all the necessary

tools for the preparation of appropriate oxide powders, sintering and

hot-pressing of ceramic bodies , flame- and/or arc-spraying equipment to

produce ceramic or metallic layers , as well as know-how and apparatus

for joining ceramic samples to a metallic base.

In order to have these electrodes built in the limited time prescribed

by the test schedule, NBS contacted two companies. The first one, Trans-Tech,

of Gaithersburg, MD, supplied the powders and ceramic bodies to our specifications, while both this company and a second one, Technetics of Milford, CT,

developed methods for making a very strong ceramic/metai bond. With the

assistance of these two companies a dozen electrode/insulator assemblies

were built and delivered to the MHD-test facility for testing in September

1976 under Argonne auspices. Details of the construction are given below.

Electrode/Insulator Assembly

Requirements: The test conditions and dimensional limitations as specified

by Argonne National Laboratory are as follows:

- two electrodes plus three insulators (cathode)

,

- two electrodes plus three insulators (anode)

,

- assembled with appropriate copper sample holders (water-cooled)

.

- length of electrodes: ^3 in.

- width of electrodes: 0.375 in.

- height of electrodes: 0.1-0. 5 in.

- insulator width: 0.05 - 0.10 in.

Two sets of the above.

- Surface temperature of electrode between 1800 and 2200 K;

- Average heat flux: o-100 W/cm^.

The test is supposed to last for 100 hrs. with "normal” start-up and

shut-down times of 30 min. each.
'

.
37

Design: The electrode material chosen for this test is the Mg-Al-Fe-spinel

(3MgAl_0 ,-lFe_0 , , now known as "MAFF-31").

2 4 3 4

The electrical conductivity as a function of temperature is given in

the 4th Quarterly Report (Fig. 1, p. 12).

The thermal conductivity was measured by Dr. J. Lambert Bates (Battelle

NW) ; data are presented in Table A .

The thermal expansion of this spinel is about 8-9 x 10 ^ at 1000 °C.

The most important design parameter is the interface temperature

between the spinel ceramic and the metal or metal alloy base. Because of

sufficient electrical conductivity of the spinel down to rather low temperature, this temperature was chosen to be about 600-700 °C. Using the

required surface temperature and heat flux, and the measured thermal conductivity of the spinel, the thickness should be of the order of 3.0 - 4.0 mm.

As a result of this small spinel thickness another problem arises:

How to attach the ceramic firmly to the metal base? Considering the

1000-1100 °C temperature difference across 3.5 mm of the spinel ceramic,

it is clear that the thermal expansion at the top surface will be much

higher than at the bottom. Hence, the strains in the ceramic layer will

be very large, resulting in a tendency for "bowing" and spalling.

To solve this problem two different approaches have been used.

a. The metal-felt pad (Technetics)

b. The cermet layer (s) (Trans-Tech)

Construction : a. Details of the electrode/insulator assembly using the metal-felt pad

are shown in Figure 17. The metal-felt is a wire-mesh material made of

an Fe-Cr-Al alloy. It is difficult to determine the thermal conductivity

of this material; the best estimate is 0.015 - 0.025 W/cm.deg. (between)

100 and 600 °C) . Hence, a thickness of 1.5 mm is required to reach a

felt-spinel interface temperature of 600 °C. In order to avoid cracking

and spalling of the spinel ceramic as much as possible, it was decided

to limit the length of the ceramic and hence to build the three-inch

electrode out of 3 one-inch "tiles".

The fabrication followed a series of steps. First, 1.5 mm thick metalfelt pads were brazed to the individual 4 mm thick copper tiles. Subsequently, a 3.5 or 4.0 mm thick layer of spinel was gradually built-up

on the metal-mesh using the flame spraying technique. After shaping

and trimming of these composites, a set of three "tiles" was brazed to

the water-cooled copper electrode holder. Final assembly of the electrode

system will be achieved by placing three insulator plates (pure dense

MgAl^O^) on each side of the electrodes and attaching them with RTV-cement.

b. The construction of the electrodes using the cermet joining technique

is quite similar. Instead of the metal-felt, the metal and the ceramic

are connected by building up several layers of different metal-ceramic

mixtures (cermets). For instance, one can flame-spray first a layer of

38

TABLE .A. Thermal Conductivity of 3MgAl20^*lFe

Temperature A

(°C) (K) (W/cm • deg)

527 800 .0317

577 850 .0319

627 900 0.0313

727 1000 .0301

827 1100 .0290

927 1200 .0279

1027 1300 .0270

1127 1400 .0262

1227 1500 .0255

1327 1600 .0246

1427 1700 .0237

1527 1800 .02233

2.963

39

ino

SLT

Q

faN6

7y

SL

7T

fa8<D

fa2H

un<n

CN

SL

-E>

•^r m

rr <r

*-S • i — fa

Design

of

spinel

electrodes

for

ANL-Reynolds

test.

40

metal powder, and then one containing about 25% of the ceramic powder.

This is followed by a 50-50 and then a 75-25 mixture. Finally, a layer

of 100% Mg-Al-Fe spinel is deposited. The total thickness of these

graded layers is small (much less than 1 mm) . On top of the cermet the

ceramic is built-up until the desired thickness is reached.

It should be understood that the particular fabrication methods

and dimensions chosen for this set of electrode/insulator assemblies

are not necessarily optimized in all respects. Many variations within

the general framework of this approach are possible. NBS is cooperating

with appropriate groups to optimize these electrodes for use in ERDA

channel development programs.

b. U-02 MHD Materials Test (Phase II) (W. R. Hosier, A.J. Armstrong, T. Negas)

During this reporting period some time was spent on several aspects

of the preparation for the joint US/USSR MHD-Materials Test (Phase II)

to be conducted in the U-02 facility in Moscow. The module (electrode and

cathode) , constructed by Westinghouse is similar to the one built for the

Phase I test (Sept. 1975). However, the electrode materials, the leadouts, the method of attachment, and the segmentation are different. The

overall length is M90 mm divided in 6 groups of 5 electrodes. Each

group (with one exception) contains 5 materials of identical chemical

composition: the last 3 groups include electrodes fabricated by two

or more suppliers (see Figure 18)

.

Mr. Hosier, who will participate in the U-02 test (Phase II) as one

of the 6 U.S. monitors, collaborated with Westinghouse staff members in

the preparation of the Work Plan and in arranging appropriate shipping

procedures

.

Powder Preparation . In cooperation with Trans-Tech Inc. sinterable

powders of conducting 3 MgAl204 : Fe 3 C>4 spinel solid solution were developed

and have been utilized by Battelle Northwest and Trans-Tech Inc. to

prepare electrodes for testing in U-02, Phase II. Grain size and phase

composition were monitored and characterized by SEM and x-ray diffraction

analysis throughout all stages of development. We are now investigating

the preparation of powders for several other Mg/Al/Fe-oxide spinel solid

solutions. Dense, sintered, ceramics are being prepared for electrical

conductivity measurements. We envision utilizing compositions lower in

iron content (hence, higher melting point) in a "graded" Mg/Al/Fe-oxide

spinel electrode provided that adequate electrical conductivity can be

demonstrated

.

Pre-test Characterization . A number of measurements , required for

pre-test characterization, were made on all samples for the U-02 test.

The results of these experiments [electrical conductivity (p. 11), structural and chemical information (Table B) ] were assembled and have been

transmitted to the USSR-side before the start of the test.
L . I m I . _
Fig. 18 U02 MODULE ELECTRODE IDENTITY

Supplier Lead-out

ASSIGNED

CURRENT

DENSITY (A /cmi

) Supplier Lead-out

41

CATHOt

?t f1

i

1 D\i 1.5 2101- ; Pt GRAuttW

CeO, -ZrO— If 1

l!02 .8 2102 1

M

1

If " 1103 .8 2103

ti 1 1 35 65

.1 " 1104 .3 2104 1

ft 50 50

75 25

73 20

II II 1105 0 2105 f i

; "

IS Inconel i205 0 2206 Inconel GRACED

II II 1207 .8 2207 ! 1 IT Ce02 —Zr02 —

II II I20S .8 2203 1? "

1 35 55

50 ’ 50

75 25

II

*

II 1209 .8 2209 II • i IT

. " " 1210 1 .5 2210 IT IT 73 20

V?- D f- 1311 1.5 2311 Pt

If II 1312 .8 2312 1

1

i

•• " 1313 .3 2313 1

TI 50CeC2 —5C

It II 1314 .8 2314 IT TI

H M !3i5 0 2315 1

! »

BNW Cu 1416 0 2415 BNW rCu

M M 1417 .8 2417 1

1

i it 75MgA!2 Qr2

Ft 1418 .3 2418 1

Pt

Trans-Tech Ft 1419 .3 2419 1

IT

bnw Cu 1420 1 .5 2420 .. Cu

BNW __ 1521 1.5- 2521 BMW Pt 75MgAl 2 O —2

- '• 1522 .8 2522 1 T

1

. "

LASL Cu 1523 .3 2523

r LASL Cu FeAl 2 0 4 -Fe 3 0

MIT Cu 1524 .3 2524 MIT ! Cu

MIT/LASii Cu 1525 1 .

2525 MIT Cu

G.R. Cu 1626 1 .2525 G.R. Cu

" " 1627 .8 2627 1

I " LaCr0 3 E

II 1623 .3 262S 9

!

*'

G.R. Ft 1629 .3 2629 G.R. ! Pt

Cu 1630 0 2530 , \>'

| Cu

PLASMA

FLOW

ANODE

- Westinghouse

BNW - Battel le Northwest

LASL - Los Alamos Scientific Lab

G.R. - General Refractories

TABLE B 42

Summary of Pre-test, U-02, Phase II Materials

Material X-ray Diffraction Analysis

(phases)

SEM/EDX

analysis

( impuri ty

)

Soinel

3MgAl

? 0

a :Fe

3 0

4

sol id-solution

(Trans-Tech Inc.)

Single-phase, wel 1 -crystal 1 ized spinel

O a=8.185A. Powders of this material

prepared in air yield a=8.160A. This

together with the fact that this

material (unlike the "in air" material

is ferromagnetic at room temperature

indicates that the Fe2 +/Fe^ + ratio is

somewhat greater than in the "in air"

spinel

None detected as major impurities

(0.5-1% level)

Spinel

3MgAl ^0^: Fe^O^

sol id-sol uti on (Battel 1 e NW)

a) Top (plasma side) portion contains

wel 1-crystal 1 ized spinel, a_=8.185A (as

Trans-Tech material).

b) Lower, more porous portion

contains at least two spinels,

a_=8.185 and a=8.170.

Cr is the major

impurity detected

(0.5-1 .0%)

Spi nel

7FeAl

2

0

4

:3Fe

3

0

4

solid-solution

(Los Alamos)

Poorly crystalline Fe/Al-oxide

SDinel with a_=8.26lA as the

dominant phase. One or more

additional spinel phases having

larger cell dimensions are also

present. Their x-ray lines are

too weak and broad for accurate

parameter determination.

Domains (<0.07 mm)

containing more porous

oxide enriched in

A1

2 03 with respect

to the denser sur- rounding spinel

matrix.

(See Fig. 23).

Spinel

7FeAl

2 0

4

:3Fe

3

0

4

solid solution

(MIT)

Well -crystal! i zed spinel with

a=8.231 A.

Contains major impurities .consisting of

K and Si (<5%) . These

are distributed

throughout the material

as porous impurity

domains (K+Si) up to

n-0.7 mm. Another set

of less dominant

impurity centers (up

to 0.05 mm) consists

of Mg and Al concentrations.

(see Figs .24, 25 )

Graded

Cerium oxidezirconia:Ta

(Westinghouse)

The composition range encompassed

by this material is described in

previous U-02 report.

Dominant impurities

(0. 5-1.0%) include

Al , Si , most of the rare earths , and probably Cs.

LaCr0 3 :Mg

(General

Refractori es

)

Well -crystal! i zed orthorhombic

perovski te 0 a = 5.478 A

b = 7.760 A

c = 5.516 A

Major impurities

(0.5-1 .0%) include

Al and probably Sc

as well as most of the

rare-earths. Minor

impurities ( <0 . 5%)

i nclude Si , Ca , and Th

LaCr0

3 :Mg

(Westi nghouse)

same as LaCr0-,:Mq

above

Major impurities Al , Si,

Sc and other rare-earths

43

The module itself was delivered in Moscow on August 23, 1976; the

U-02 test (Phase II) is expected to begin on September 20, 1976 and will

last for at least 100 hours. After the test, the module is expected to

be brought back to NBS around October 1, 1976; it will be disassembled

during the week of October 4 in the presence of several Soviet representatives.

c. U-25 Bypass Channel (H. P. R. Frederikse , S. J. Schneider)

During the Spring meeting of the US/USSR MHD-Steering Committee in

June 1976, it was suggested that future tests of electrode materials and

generator walls could be conducted in a test channel built as a bypass of

the large U-25 generator duct. This bypass channel will be constructed

to fit inside a superconducting magnet to be delivered to the USSR in 1977.

The magnet will provide a field of 5 Tesla. As far as the channel is

concerned , it is intended to be operated with maintenance material added

to the fuel, and a mass flow appreciably higher than that of the U-02.

The channel would be used both for plasma experiments and wall materials

tests. The implementation of these plans will be discussed at the Moscow

meeting of the MHD-Materials Group in October 1976.

d. Soviet Test at UTSI (US/USSR MHD-Materials Testing)

Two Soviet electrode materials will be tested in the UTSI MHD facility

in early October. This test will be run with coal slag additives. A. J.

Armstrong will participate in this test. The two materials are SiC (+ Ti)

and Cr-LaCrO^. The first batch of samples has been received at NBS. All

but 4 of each (for X-ray measurements) have been forwarded to Dr. Lloyd

Crawford at UTSI.

e. AVCO Test (August 1976) (T. Negas , W. Hosier, E. Farabaugh, L. Cook)

In September 1976 we received two anode electrodes from AVCO Corp.

for analysis . Stan Petty conducted an MHD test utilizating cooled

rectangular copper bars onto which a metal alloy mesh (1.5mm thickness)

produced by Technetics Corp. was brazed. The composition of the alloy

wire is considered proprietary but contains Fe and Cr as dominant

constituents with A1 and Si as minor components. One anode mesh was

plasma sprayed with a conducting spinel solid solution having the composition 3MgAl2 04 :Fe 304 . Spravable powders of this material were prepared

by Trans-Tech Inc. The thickness of the sprayed spinel layer was ^l.Omm.

_:.e second anode was sprayed (vi.Omm) with 80ZrC>2/20Y 0 from powders of unknown

origin. These anodes first were equilibrated in the AVCO test channel

with AVCO slag before seed was introduced and power drawn. Test duration

in the slag-seed-power mode was 5 hr. Electrode design was by AVCO although

the conducting spinel was suggested by NBS.

The two anode bars were "potted" with epoxy and several slices perpendicular to the bar length were sawed from each and lightly polished for

analysis by SEM/EDX* Initially, there was some question concerning the

direction of plasma flow with respect to anode orientation as received.

However, during analysis it became clear that one side of each electrode

consistently revealed maximum degradation, particularly of the metal mesh.

This side is assumed to be the upstream portion of the anodes where current

tends to be concentrated by J x B

.

44

It mav be recalled that from the analysis of previous AVCO anodes the

nature of the degradation of inconel electrodes could not be deduced

definitively as very little anode metal remained after long duration (100

hr.) testing [1]. The results from the present test suggest that localized

arcing and heating primarily on the anode upstream side lead to the

degradation of anode metal-ceramic components. Both anodes, regardless

of the initial, intervening ceramic oxide layer between wire mesh and

slag, reveal similar phenomena. The sprayed ceramic layers tend to be

discontinuous and crystalline (cold) slag overlies either these layers or

is directly on the wire mesh. On the upstream side where slag mainly

overlies metal mesh the alloy is degraded (oxidized) but not extensively

for this short test. On the downstream side the alloy remains relatively

unaltered whether covered by ceramic plus slag or mainly crystalline slag

without ceramic. These observations suggest that slag dissolves the ceramic

oxides (particularly the Zr02 ~based material) with or without high current

activity. However, degradation of the wire mesh is apparent only in the

upstream high current density region.

Figure 19a is typical area (150X) revealing the wire mesh-ceramic oxide

(80Zr09/20Y 7 03) -slag interfaces. Figure 19b, a corresponding map for Zr

clearly shows what remains of the sprayed zirconia layer. Wire under the

remaining zirconia (at right) remains in relatively good condition. Where

zirconia is nearly absent (at left) the area has been invaded by slag

components (K and Si) and the wire was subjected to oxidation. This is

illustrated by Figs. 19c (K map) , 19d (Cr map) and 19e (Fe map) . Notice

(Figs. 19d and 19e) that for this alloy composition the Cr component is

particularly sensitive to oxidation.

Figure 20a (100X) is a typical area showing the wire mesh-ceramic oxide

( 3MgAl204 :Fe 304 spinel) -slag interfaces. Here the remaining spinel has a maximum thickness of 0.2-0. 3mm and the underlying wire is in relatively

good condition. Figure 20b (Mg map) and Figure 20c (Si map) reveal the

spinel-slag interface. Figure 21 ( 190X) is an adjacent upstream area of

the anode. Here an interface between two unlike materials exists along

the vesicular zone running north-south about midway in the figure. Figure

22a is a magnification (500X) of a portion of this interface. Included is

a wire remnant also shown in Figure 21. Figure 22b (Si map) reveals the

position of the lower portion of the slag layer. Figures 22c (Fe map) and

22d (Cr map) show that under the slag is a zone of oxidation enriched with

Fe and particularly Cr , the dominant wire components. The spinel phase is

practically absent. Notice, as before, that the wire remnant is depleted

in Cr.

Several conclusions can be drawn from these data as well as other

extensive photographic evidence available but not presented. These include:

1. a rather sharp interface can and does exist between AVCO slag

(crystalline or glassy) and sprayed ceramic suggesting compatibility

after considerable recession of the latter;

this ceramic-slag zone, however, is discontinuous and practically

absent at the anode upstream side where the wire mesh is subjected

to oxidation, particularly for the zirconia-mesh anode; and

2 .
'
45

3. this correlation indicates that localized, excessive heating and

high current density associated with arcing result in,

a) solution/incorporation of the sprayed ceramic into/by the slag

phase , and

b) oxidation of the underlying wire mesh, probably without melting

of this particular alloy.

It may be supposed that conducting ceramic oxides thicker than ^lmm

may be more beneficial to "protect" underlying metal anodes. However, if

arcing proceeds through slag-ceramic-metal interfaces, even greater

thicknesses should be subjected to similar destruction. In the arcing

mode, therefore, the situation boils down ultimately to slag on metal

whereupon the properties of the metal become overwhelmingly important.

It should be noted, however, that a comparison of the sprayed

zirconia and spinel anodes reveals that, for this short test, the

latter (better conducting) material appears to be beneficial in

terms of "spreading out" current distribution. The zirconia anode

shows evidence for current concentration only on the upstream side.

We also received from AVCO unoriented samples of "slag" from tested

(20 hr) cathodes consisting initially of metal (?) -hercynite (Fe-Al

oxide spinel solid solution of unknown bulk composition) . The material

proved to be nearly identical with slag from cathodes analyzed in

reference [1] . Hercynite -type spinel was not found. Two tested

cathodes consisting of the above materials also were received at the

end of this reporting period. Analysis is in progress.

REFERENCE

[1] Analysis of AVCO Mark VI C Channel Electrodes After Long Duration

Test, W. R. Hosier, T. Negas and S. W. Petty, Proceedings 15th

Symposium on Engineering Aspects of Magnetohydrodynamics, Philadelphia,

PA, May 24-26, 1976.
'
Fig. 19a. Anode (150X, upstream area) consisting of metal alloy (Fe, Cr)

mesh (bottom) , the remains of a Q07.rO 20Y ^0^ sprayed layer

and slag (top)

Fig. 19b. Elemental map for Zr of area in Figure 19a.

Fig. 19d . Elemental map for Cr or area in Figure 19a.

Fig. 19e . Elemental map for Fe of area in Figure 19a.

48

Fig. 20a. Anode (100X) consisting of metal alloy mesh (bottom) , the remains

of a conducting spinel layer (see text) and slag (top)

Fig. 20b. Elemental map for Mg of area in Fig. 20a showing position of slag layer.

Fig. 20c Elemental map for Si of area in Fig. 20a showing the slag layer

Fig. 21. Anode (190X, upstream area) adjacent to area in Fig. 20a

showing wire, oxidized wire and slag.
J
50

Fig. 22b. Elemental map for Si of area in Fig. 22a.

51

Fig. 22c. Elemental map for Fe or area in Fig. 22a.

Fig. 22d. Elemental map for Cr of area in Fig. 22a.
J
52

Fig. 23 . Typical Al-rich domain

(1000X) in Fe/Al-spinel prepared

at Los Alamos

.

Fig. 24 . Typical Mg/Al-rich

domain (2000X) in Fe/Al-spinel

prepared by MIT.

Fig. 25 . Typical K/Si-rich domain

(1000X) in Fe/Al-spinel prepared

by MIT.

53

5EM Examination of Corhart X-317 Ceramics (Magnesia Spinels) Tested at

FluiDvne Eng. Corp. (E. N. Farabaugh and D. B. Minor)

One of the most promising preheater materials is the Corhart X-317

spinel. Specimens of the materials tested under different experiment

conditions have been examined. The object of these studies during this

reporting period has been to determine the depth of penetration of

different elements into the spinel, to identify elements in the reacted

area, and to determine which of the two phase MgO and MgA^Ot*, present

in the X-317 materials, is being predominately reacted in the near

surface volumes

.

Specimen FN-27 (Cohart X-317) was run at 1800°K for 30 hours in the

presence of ash and seed. SEM mapping and EDX work has shown that the

reacted layer is ^0 . 4mm thick. Its microstructure has more large pores

than in the bulk of the specimen. Thus, one can roughly determine the

limits of the reacted layer by noting the position where the change in

microstructure is seen. This is shown in Fig. 26 where the arrows indicate

the extent of the reacted layer.

Elemental concentrations of Mg, Al, Si, Ca, K, Cr, and Fe were

found in the reacted layer. The concentration of Mg increased as one

moved from the surface through the reacted layer into the bulk of the

material where it leveled off. Al concentrations varied but essentially stayed

within fixed limits over the same region. The x-ray peaks for Si, Ca,

K, Cr , and Fe all decreased as the layer was traversed, falling to near

background at distances of twice the layer thickness (^.8mm).

To address the problem of which phase is reacting predominately , x—

ray mapping was done on a portion of the reacted layer shown in Fig. 27a.

Fig. 27b is an Al map, 27c is a Ca map, 27d is a Si map and 27 e, a map of

the same area. Note that in 27b we see areas with no Al concentrations.

These areas are filled in with Ca and Si as shown in 27c and 27d. In 27e

there is no large concentration of Mg associated with these areas. A

possible explanation of these features is that the Al-free areas in 27b

were originally MgO grains. The area diameter, r'j 4 0 tm, is about the right

size for MgO grains found in the untested material. Upon testing, Si,

Ca , K, and Mg may react so as to attack these MgO grains forming other

phases in their place. This would explain the lower concentration of Mg

in the reacted layer relative to the bulk, as well as the appearance or

the elemental maps.

To check this idea further, part of the specimen in the bulk (4-5

mm from the surface) was examined. The region is shown in Fig.2 8a , where

we can see the outline of grains in the surface. Fig. 28b is an Al

map of the same area showing that the grains outlined in 28a coincide

with the Al-free regions in 28b. Fig. 28c is an Mg map of the same region.

The high Mg concentration coincides with the grains of 28a. Thus we can
J

I
54

conclude the grains are MgO. This type structure was reacted in the gas

streams to produce that shown in Fig .27a during its test. These results

suggest that the MgO is preferentially reacted compared to MgA^Oi* in the

test. MgAl?Ou. also is reacted, but not to the extent that the MgO

grains are.

X-ray data taken by A. Perloff tend to support this explanation.

Powder made from the reacted layer showed less MgO relative to MgA^Oi* than

powder made from the bulk of the specimen.

A micrograph taken from specimen FN-65 (Corhart X-317) tested for 30

hrs . with ash and seed and cycled between 1800°-1600°K is shown in Fig. 29a

The depth of penetration was 'v.lmm. Maps of the same region are

shown in 29b for Mg, 29c for Al,29d for Si. Again the Mg deficient

reacted layer is seen, suggesting the preferential reaction of MgO with

gas stream components. Fig. 29c shows a more uniform distribution of A1

in the reacted layer than in the bulk, suggesting the MgAl 2©4 has also

reacted. 29d shows the higher concentration of Si in the layer. Elemental

constituent of the layer were Si, K, Ca, Mg, Al, Cr, and Fe.

Fig. 30 is a micrograph of FN-70 (Corhart X-317) tested with seed

but no ash for 30 hrs. and cycled between 1800°-1600°K. The depth of

penetration here is ^.2 mm. Elemental maps gave results much the same

as seen before for the two previously discussed specimens, A Mg deficient layer which contained more Si and K than the bulk. There was some

Fe present in the spectra taken from the layer but not as much as was

present in the spectra of FN-65. Since there was no ash in this run, it

is suggested that Fe must have come from the Mg-Cr bricks used to line

the test channel. Elements found in the layer were again Si, K, Ca, Cr,

Fe, Mg, and Al.

Virgin X-317 was also examined. Maps and x-ray spectra were taken

of the specimen. The results showed that the microstructure of the

virgin material was identical to that of the bulk of the tested material.

Grains of MgO were located and identified which were much the same as

shown in Fig. 28a. X-ray spectra showed the presence of Mg, Al, and Ca.

Thus, the Ca present in the tested materials was present in the virgin

material and not necessarily a result of the testing.

Fig. 26 . 100X secondary electron micrograph of Corhart X-317 specimen

(FN-27) showing the depth of the reacted layer.

Fig. 27a. 500X secondary electron

of the reacted layer of

micrograph taken in the polished surface

FN-27

.
a
Fig. 27b. Al area map of the same area.

Fig. 27c. Ca area map of the same area.

Fig. 27d. Si area map of the same area.

Fig. 27e. Mg area map of the same area.
I
58

Fig. 28a. 500X backscattered micrograph taken in the bulk of FN-27

showing grains in the MgAl^O^ matrix.

Fig. 28b- Al area map of the same area shown in Fig. 28a.

Fig. 28c. Mg area map of the same area shown in Fig. 28a.

60

Fig. 29b. Mg area map of the same area as in Fig. 29a.

61

Fig. 29c. Al area map of the same area as in Fig. 29a.

Fig.29d. Si area map of the same area as in Fig. 29a.
j
62

Fig. 30. 200X secondary electron micrograph of the reacted layer of

FN-70

.
j
63

Task L. ASSESSMENT OF STEAM PLANT COMPONENTS (J. R. Cuthill)

This is the 5th quarterly progress report since the initiation of

Task L. The "Baseline Plant System Design Description" (1) became

available at the beginning of the quarter. Figure 31 is a redrawn

block diagram of the system from the combustor through the; steam reheater,

giving the flue gas temperature at each station and the specified

steam temperature and pressure in the steam superheater and reheater.

It will be noted that the flue gas temperature is higher in the steam

superheater and reheater than in the previous STD calculation, Fig. 32

Another, and possibly more significant difference is in the composition

of the flue gas through the steam superheater and reheater, between

the previous STD calculation and the Baseline Plant Calculation. The

potassium seed going through the superheater and reheater is predominantly K2 CO 3 and K2 SO4 approximately in the ratio of 2:1 in the

previous STD calculation, but is predominantly in the form of sulfate

in the Baseline Plant calculation. The latter situation may be a step

in the right direction from the point of view of hot-corrosion attack

judging from a remark by Stringer in a review paper (2) that an

acceleration of oxidation had been observed in a salt-coated test when

Na^CO^ was substituted for Na^SO^.

Hot-corrosion Attack . If the mechanism of hot corrosion attack

follows the pattern in marine gas turbines involving sodium compounds,

the severity of the K^SO^ attack will depend strongly on the amount of

KCl present, which fortunately is small but not insignificant, about

1% of the amount of K2 SO4 present. The effect of the chloride ion

content will be discussed in more detail later.

In the Baseline Plant Design, 304H stainless steel [ ASTM Spec.

A(213) 304H],ana 2-1/4 Cr - 1 Mo low alloy steel [ASTM Spec. A(213)T22]

(3), is designated for the steam superheater and reheater coils respectively.

In view of these steels being specified for the superheater and reheater

units it would seem particularly pertinent to summarize the hot corrosion

data on these alloys.

Crolov 2-1/4 which complies with SA(213)T22, was included in the

simulated MED corrosion test (4) by the Bureau of Mines at Bruceton (now

PERC) . (4)

.

The Croloy 2-1/4 was attacked at a wall temperature of 800 °F

and a gas temperature of only 1800 °F. They concluded that Croloy 2-1/4

could fulfill "short-term" use at these conditions.

An equivalent steel was included in the simulated MHD hot corrosion

test conducted in the Central Electricity Research Laboratories, England. (5)

They concluded that 2-1/4 Cr , 1 Mo could be used up to 630 K (675 °F) . Note

(Fig. 31 ) that the wall temperature would probably be in this vicinity.

An 18 Cr-12 Ni , 1 Nb stainless steel was included in the test program.

It showed severe pitting above 725 K rendering it unsafe to use above this

temperature. Unfortunately, an 18-8 was not included in the program.

Both the PERC and CERL studies are summarized in the ECAS , materials

section, ( 6 ) but the hot corrosion aspect does not seem to have been

taken into account (p. 3-208) in Table 3.10, in which the 304 stainless

steel and the 2-1/4 Cr , 1 Mo steel is recommended for the superheater

and reheater, respectively.
,

j.
64

On the basis of the data available, the 2-1/4 Cr , 1 Mo low alloy

steel would not appear to be a viable choice for the reheater, and

probably neither is the 304H stainless steel in the case of the

superheater unit. Of course, it is understood that such a plant will

not be built for a number of years. During the intervening time a

continuous reappraisal should be made of the materials selected for

the respective components as successively better data becomes

available on the actual service conditions to be encountered. For

instance, the amount of KC1 present with the sulfate appears to be

particularly critical.

Schultz and Halsiger (7) discuss results from two different INCO

research laboratories using the same "crucible" hot corrosion screening

test in their alloy development programs except one used a 10% NaCl90% Na 2 SC>4 molten salt bath and the other used a 25-75% mixture. The

latter was much more severe as will be seen by comparing the corrosion

rates of IN-738 and UDIMET-500 in Fig. 33 with those in Fig. 34.

However, even more spectacular, was the finding that hot corrosion

could not be induced in IN-713C in any reasonable length of time when

the NaCl was omitted entirely. Also, Stringer (2) warns against

increasing the chloride ion content to accelerate corrosion attack

because the hot corrosion dependence on chloride concentration has

not been established making extrapolation difficult.

For this reason, experimental data on the hot corrosion attack

under the specific anticipated service condition needs to be obtained.

A test rig that duplicates as nearly as possible the service conditions

is unquestionably the straight forward approach for obtaining meaningful data. However, because such tests are expensive and time consuming,

in the gas turbine case as well as the MHD case, there have been a number

of types of "crucible" tests exploited to try to obtain meaningful

correlations with actual service. Schultz and Halsiger (7) have

recently described a version called the Salt Shower Test being used at

the INCO European Research and Development Center in Birmingham, England.

In this test the specimen is mounted on a slowly rotating spindle while

maintained at 900 °C or other chosen temperatures in still air. The salt

mixture, in the form of dry powder (8)

,

is "showered" down on the specimen

either continuously or for 5 minutes every hour. [Fig.35(9)]. This dry powder melts

when it hits the hot specimen, running down over it and coating it. In

addition to its simplicity and flexibility, this test is reported to

overcome the shortcomings of an immersion test in molten salt (which

they used previously ) in that it will reliably cover cobalt base, as well as

nickel base alloys over a wide range of chromium content and a wide

spread in resistance to hot corrosion attack.

The compositions of the three wrought alloys showing the exceptional

low weight loss in Fig. 34 , i.e., M313, IN-587, and IN-597, (IN-939 is

a cast alloy) are given in Table 1. IN-597 is reported to be comparable

in strength properties with wrought Udimet-500. Some pertinent mechanical

properties of Udimet-500 (10) are tabulated in Table 2.
$
65

Future Plans

The immediate emphasis will be on summarizing the hot corrosion data,

particularly potassium compounds such as, K^SO^, K^CO^, KC1 and KOH, on

the structural alloys being considered. While there is a wealth of data

on the hot corrosion attack by sodium compounds, from the marine gas

turbine literature , there is comparatively little on potassium compounds

but not insignificant. This' data on attack by potassium compounds will

be summarized.

Another compilation that appears worthwhile is a listing of alloys,

considered for heat exchanger tubes, in the order of their resistance

to hot-corrosion attack. This ranking will, of necessity, have to be

largely on the basis of the marine gas turbine work, because of the

limitations on the availability of data.

References

1. Baseline Plant System Design Description, June 30, 1976, by Gilbert

Associates and STD Research Corp.

2. John Stringer, Hot Corrosion in Gas Turbine, pp. 79-101, Corrosion

Problems in Energy Conversion and Generation, C. S. Tedmon, Jr., Ed.

1974, the Electrochemical Society, Princeton, NJ.

3. 1976 Annual Book of ASTM STandards , Part 1, American Society for

Testing and Materials, Philadelphia, PA.

4. D. Bienstock, R. J. Demski , and R. C. Carey, Corrosion of HeatExchange Tubes in a Simulated Coal-Fired MHD System, Trans, of ASME

,

J. of Engineering for Power, 93_, A, #2, pp. 249-256.

5. Open Cycle MHD Power Generation, p. 721, J. B. Heywood and G. J. Womack,

eds., Pergamon Press, New York.

6. Energy Conversion Alternative Study, Westinghouse , Phase I, Final Report,

Vol. II - Materials Considerations, NASA CR-134941, NTIS , Springfield, VA.

22161

7. J. W. Schultz and W. R. Halsiger, Corrosion Resistant Nickel-Base

Alloys for Gas Turbines, Metals Engineering Quarterly, 16_, 3, pp . 15-24

(1976) .

3. Private Communication.

9. ?. L. Norman and J. D. Harston, An Evaluation of Commercial HighChromium Nickel-base Alloys , Deposition and Corrosion in Gas

Turbines, A. B. Hart and A. J. B. Cutler, eds. Applied Science

Publishers, Ltd., London, 1972.

10. Aerospace Structure Metals Handbook, Vol. 4, Metal Properties Data

Center, 3elfar Stulen, Inc., 13919 West Bay Shore Drive, Traverse

City, Michigan, 49684.
. I
Table

1.

Wrought

Alloy

Composition s

66

a

fao

n ino

r~O

v

mi

in cm i i—i ws

hHQ3

fa

Oiao

fa

fa ooo

CN I I

cn

o

o • — • • • • rH CN ro ro

CN

m

O

—i • • • • < O rH rfa ro O o O 0 o o &

CJ 1 CN CN

r

—1. rH rH rfa rH

•fa 03 03 td 03 z oi A fa fa C ln in o

fao CO rr <y>

u

m CN CN rHoo

in

p-~ 1 Cfa bo (Ti H

m

n

n 2

rH i i H

cn z Z fa % M M D

fa0

in0

•fa fa

fa0

fa0

fa fa

faS'

<D

fa facn I

fan30

fa

<d

cn0

•H

fa

fa<u

fa0

fafa

CN

0)

rfa 9

E-<

n

o

mo i—i

ro fa)

rfa rH pH •• •••••• •1•

•rH

•fa cn

Cn fa

IT • *.

*. cn

r-> fa • fa fa

cn • C •rH

o 0 • cn o

fa •H fa

o

fa cn i—i

cn fa -

_r; fa cu

*. fa

o

fa fa cn fa 3 fa fi fa

fa cn CU cd Oj c fa 3 CU fa cu

fa fa cn 0

fa 0 , cn <U fa

fa

r—1 V

fa TO •fa i—i

cn dP

O cu

r-* rfa

o •fa 3 • rH >1 &H o

Figure

31.

Baseline

Reference

Design

for

an

Open

Cycle

MUD

Power

Plant

FEED WATER 67

TO ECONOMIZES AND STACK
' ,v 'i m
WAT

EK

FLOW

__

.

..

,

STACK

68

X

Fig.

32.

STD

Research

Corp.

state

point

temperatures

with

seed

condenser

after

the

steam

plant

69

Fig.

Fig

.

33. Comparison of superalloy performance in

90 pet Na SO -10 pet NaCl crucible test at

927 °C.

34. Comparison of superalloy performance in

75 pet Na

2

S0

4

~25 pet NaCl salt shower test

at 900 °C7

70

Fig. 35 . Schematic diagram showing essential features of the salt shower rig.
)
76-1151 CANCELLED