FE-17A9-8

Dist. Category UC-90C

MATERIALS RESEARCH FOR CLEAN UTILIZATION OF COAL

QUARTERLY PROGRESS REPORT

July 1976

Samuel J. Schneider

Project Manager

Institute for Materials Research

National Bureau of Standards

Washington, D. C. 20234

PREPARED FOR THE UNITED STATES

ENERGY RESEARCH AND DEVELOPMENT ADMINISTRATION

Under Contract No. 14-32-0001-1749

"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

subcontractor, or their employees, makes any warranty, express or implied, or assumes any legal liability or responsibility for the

accuracy, completeness, or usefulness of any information, apparatus,

product or process disclosed, or represents that its use would not

infringe privately owned rights."
/

Scnneider, S. J., Materials research for

clean utilization of coal. Quarterly

progress Report July 1976, Energy Research

and Development Administration FE-17^9~8,

73 pages (Available from the National

Technical Information Services, Springfield,

Va 1976).

313

ms*
TABLE OF CONTENTS

I.

II.

III.

OBJECTIVE AND SCOPE OF WORK

SUMMARY OF PROGRESS TO DATE

/ DETAILED DESCRIPTION OF TECHNICAL PROGRESS

1. Durability and Reliability of Materials of Construction. . . .

A. Metals

a. Corrosion

b. Erosive Wear

B. Ceramics

a. Deformation and Fracture

b. Erosive Wear

2. Chemical Degradation of Ceramics

A. Gas Phase Reactions

3. Consultation and Proposal Review

4. Failure Analysis

PAGE

3

3

3

3

18

38

38

46

58

58

60

70

A. Failure Avoidance Program 70

I. OJBECTIVE AND SCOPE OF WORK

Coal Gasification processes require the handling and containment of

corrosive gases and liquids at high temperature and pressures, and also

the handling of flowing coal particles in this environment. These severe environments cause materials failures which inhibit successful and longtime operation of the gasification systems. The project will entail

investigations on the wear, corrosion, chemical degradation, fracture,

and deformation processes which lead to the breakdown of metals and

ceramics currently being utilized in pilot plants (and suffering failures)

Studies will also be carried out on new candidate materials considered for

improved performance. Special emphasis will be devoted to the development

of test methods, especially short-time procedures, to evaluate the durability of materials in the gasification environments. These methods will

focus on wear, impact erosion, stress corrosion, strength, deformation,

slow crack growth and chemical degradation of refractories. Failure analysis

of gasifier components, both metal and ceramic will be conducted as necessary.

Failure reports from gasifiers will be compiled, abstracted and recommendations made to ERDA as to the appropriate action to be taken. All studies

will be correlated with field inspections and failure analysis on in-service

materials. Active consultation and proposal review services to ERDA and

associated contractors will be provided.

II. SUMMARY OF PROGRESS TO DATE

The summary of progress is contained in the individual task progress and

plans - Section III. See Performance Chart following.

1
PERFORMANCE AS OF July 1, 1976' OCT.

1975

JAN. 1976 APRIL

1976 JULY1976

Metal Corrosion

Complete evaluation of the constant strain rate technique with a system of known susceptibility to SCC.

Begin evaluation of the constant strain rate technique with materials

and environments of relevance to coal conversion technology.

Metal Wear

Characterize surface films and coatings on test specimens.

Conduct study of sulfide environment effects.

Examine effects of high pressure on erosion/corrosion.

Obtain and examine in-service specimens.

Develop and quantify erosion test method and equipment.

Ceramic Deformation and Fracture

Design controlled atmosphere apparatus for mechanical property

measurements in reactive gases at elevated temperatures and elevated pressures.

Measure strength of calcium-aluminate-bonded refractory castable after exposure to elevated-pressure steam at temperatures from 100 to 510°C.

Identify hydration phases of calcium aluminate after exposure to elevated-pressure steam at temperatures from 100 to 510°C.

Measure strength in candidate refractories as a function of loading

rate to failure at elevated temperatures without reactive environments.

Construct controlled atmosphere apparatus for mechanical property measurements in reactive gases at elevated temperatures and elevated pressures.

Collect compressive creep data of candidate refractories at elevated

temperatures without reactive environments.

Correlate s low-crack-growth , strength, and creep behavior in the absence of reactive environments.

Collect strength data in reactive gases at elevated temperatures and elevated pressures.

Erosive Wear Ceramics Erosive wear of calcium aluminate bonded alumina-dependence of wear on temperature, velocity and impingement angle.

Erosive wear of other generic castable refractories.

Design and construction of improved apparatus to investigate wear

in controlled environments.

Erosive wear study of refractories that have been exposed to high pressure steam.

Chemical Degradation of Ceramics Assessment of pertinent published phase equilibrium diagrams.

Development of rough screening tests.

Chemical attack of slag

a. commercial refractories

b. controlled laboratory specimens

Assessment of gas-refractory reactions.

Equilibrium studies of the CaO-Al 0 -H 0 system in the P and T regions

of interest in coal gasifiers.

In conjunction with mechanical properties section correlate chemical

stability, strength with phase assemblage.

Determine phase assemblage for laboratory specimens.

Obtain in service specimens and examine phase assemblage.

Proposal Evaluation and Consultation

Participate in ERDA program reviews

Failure Analysis - Metals and Ceramics

Abstract failure analysis reports made to ERDA and recommend appropriate

action, as needed.

Conduct failure analysis on components from coal conversion plants

as needed.

Solid bar refers to percent completed.

2
III. DETAILED DESCRIPTION OF TECHNICAL PROGRESS

I I. Durability and Reliability of Materials of Construction

A. Metals

a) Corrosion (G. M. Ugiansky and C. E. Joh'nson, 312.04)

Progress : Work has continued with the testing of 310 stainless steel. The

constant strain rate technique has been shown to be useful for detecting

embrittlement of 310 SS at 1000°F which is dependent on strain rate but

independent of environment. Lower ductility was exhibited for 310 SS at

a strain rate of approximately 1 x 10-6 S~1 in inert atmospheres (He, Ar) as well as H^S and H 2 S/H

2O gases. The specimens tested previously in the inert

atmospheres had some oxide on the surface even though the gas was passed

through a liquid nitrogen trap to remove water vapor. The oxide also could

have been formed from the low concentration (1.5 to 5 pnm) of oxygen in the

gas. Several more 310 SS specimens were tested in inert atmospheres using an inert gas purifier to remove oxygen. Liquid nitrogen cold traps were also

used to remove the water vapor. Fig. 1 is the stress-strain curve for a 310 SS specimen tested at 1000°F at the slow strain rate of 8.43 • 10~ 7 S"1

in a helium atmosphere using the inert gas purifier. The fracture of this

specimen was brittle with the accompanying low ductility and secondary cracking. A light gray oxide was still present on the surface of the specimen.

Fig. 2 is a stress-strain curve of a 310 SS specimen tested in argon using

the inert gas purifier at a slow strain rate of 8.4 • 10*7 S“1 at 1000°F.

The fracture was also brittle with the accompanying low ductility and secondary

cracking. A light gray oxide was still present on the surface of the specimen.

Another 310 SS specimen was tested in ultra-pure helium at a static pressure

(no flow) with liquid nitrogren cold traps at both the entrance and exit to the

test cell. Fig. 3 is the stress-strain curve for this specimen which again

illustrates the brittle fracture and low ductility at the slow strain rate

of 8.4 • 10*7 s-1 at 1000°F. A light gray oxide was again present on the surface

of the specimen. To date we have been unable to obtain a tested specimen that

was free of any surface oxide when using various inert atmospheres. Another

310 SS specimen was tested this time in vacuum. The vacuum was attained by the

in-house vacuum line (726 mm Hg). Fig. 4 is the stress-strain curve for this

specimen tested at 8.4 • 10~7 S"1 at 1000°F. Again the fracture was brittle

with low ductility and secondary cracking. A light blue oxide was present

on the surface of the tested specimen. Because of the color, the oxide

appears to be thinner than the ones obtained from the inert atmospheres.

Since we have been unable to obtain an oxide free specimen when tested in an inert atmosphere or vacuum, we proceeded to test 310 SS at different strain

rates and temperature assuming that the oxide on the surface of the specimens

was not part of the mechanism that produced brittle fracture. Fig. 5 is a stress-strain curve of a 310 SS specimen that was tested at a faster strain

rate of 1.3 * 10“ 4 S" 1 at 1000°F in a helium atmosphere where the gas was passed

through the inert gas purifier and a liquid nitrogen trap. The fracture

was ductile with no secondary cracking. The stress-strain curves in Fig. 6

show the effect of strain rate on two 310 SS specimens tested at 1000°F in

dry helium. These curves could also be interpreted as the effect of time at

3
temperature on the properties of stress and strain. The embrittlement

phenomenon seen when testing at the low strain rate in either inert atmosphere

or H

2 S is believed to be due to strain induced diffusion of interstitial

impurities to grain-boundaries causing a brittle phase to form. Sensitization

(formation of chromium carbide in the grain boundary thereby leaving a chromium

depleted region) is known to occur in 310 SS in the temperature range from 800

to 1650°F. With this in mind, a 310 SS specimen was tested at 700°F at a slow strain rate of 8.43 • 10” 7 S

_

1 in dry helium using the gas purifier. The

stress-strain curve for this specimen is shown in Fig. 7. The fracture was ductile with no secondary cracking. Fig. 8 shows the effect of temperature

on the stress-strain curves of 310 SS when tested at a slow strain rate of

8.43 • 10“ 7 S

-1 in a dry helium atmosphere. The curves do seem to show evidence

of embrittlement within the temperature range at which sensitization occurs,

however, when tested below this range the material remains ductile. The results

of a compositional analysis of the 310 SS were as follows:

Carbon 0.08%

manganese 1 IFv1 • / 0

1 nospnorus n n°7

n n° 1 ou 1 t u r

oi n a° 1 "i c 0 n U • ‘HO

1 O Cf. M i c«e

1 y . D D OA Q7 unrorn 1 um - LH. OJ

1 n mo lyDaenum . 1 u uopper 1 1 7 Iron —- Bal

With the amount of carbon present, chromium carbide could readily be formed

at the test temperature of 1000°F. The carbide forms very slowly only when the

carbon content is on the order of 0.03% or less [1],

Other evidence that sensitization is occurring_at the 1000°F test temperature

and not at 700°F is shown in Figs. 9 and 10. Figs. 9A and B are photomicrographs of a 310 SS specimen as received and after testing at 1000°F at the

slow strain rate of 8.43 • 10" 7 S"1 (brittle fracture), respectively. Fig. 9A

shows the grain boundaries to be partially etched indicating some sensitization

or carbide precipitation. Fig. 9B shows that the grain boundaries are completely

etched indicating that further sensitization has occurred during testing.

Figs. 10A and B are photomicrographs of a 310 SS specimen as received and

after testing at 700°F at the slow strain rate of 8.43 • 10“ 7 S

_

1 (ductile

fracture), respectively. Fig. 10A shows the grain boundaries to be partially

etched indicating some sensitization or carbide precipitation. Fig. 10B shows

that the grain boundaries are not more deeply etched than in Fig. 10A

indicating little or no further sensitization or precipitation occurred

while testing at 700°F. (The same etching time was used for all specimens).

The brittle fracture which could be caused by the carbide precipitation only

occurs at very slow strain rates. This point was made in an earlier report

when a 310 SS specimen was held at 1000°F for 72 hours under a no load condition

and then tested at a fast strain rate at 1000°F. The result was a ductile

fracture with 34% elongation. If the temperature or time at temperature causes

the sensitization which in turn causes the embrittlement, then the above

test would have shown a brittle fracture. The result shows that a necessary

condition for brittle fracture is a slow strain rate at temperature indicating

4
either a complicated mechanism or more than one mechanism of brittle failure.

A literature search was started on some work done in the 1950's and early

1960's regarding the strain rate dependent embrittlement of austenitic alloys.

The literature search has shown that there may be several mechanisms that

cause the brittle failure. Phase transformation in the alloy could be occurrino

when it is being plastically deformed during testing. Slow strain rate

embrittlement or strain aging embrittlement (a facet of slow strain rate

embrittlement) could be a mechanism for failure. The straining and aging

occurs during the test when the strain rate and temperature are appropriate to

allow sufficient mobility to a particular interstitial element involved. The

migration of the interstitial element to grain boundaries, for example, can cause a brittle phase to form. Some of the interstitial alloying elements

which have an influence on the mechanical behavior of alloys are carbon,

nitrogen, oxygen, and boron. Of the literature searched, all the authors

[2, 3, 4, 5] agree that the delayed brittle failure can only be detected by

slow strain rate tests.

Plans : Test 347 and 310 S stainless steel and compare data with the 310 stainless steel data. The stabilized 347 SS is recommended for service where intermittent heating and cooling occurs between 800 and 1650°F to avoid intergranular

embrittlement, whereas, the 310 stainless steel is recommended for service

at constant elevated temperatures . The 310 S stainless steel which is low in

carbon to reduce the precipitation of intergranular carbides can be used in

the 800 to 1650°F zone for several hours but is not satisfactory for longtime service in the sensitizing temperature range because it is not completely

immune to the formation of carbides which are detrimental to the corrosion

resistance of the alloy. The literature search for data regarding strain-rate

dependent embrittlement of austenitic alloys will continue.

Hi qh 1 i ghts : Early in the year, smooth round brass specimens were tested in

water vapor saturated ammonia and were found to be susceptible to stress

corrosion cracking in the gaseous atmosphere at a strain rate that was roughly

within an order of magnitude of what was observed for brass in an aqueous

medium. It was found that elongation at failure and reduction in area were

better indicators of susceptibility to stress corrosion cracking than

ultimate tensile strength. The ultimate tensile strength depended on the

state of the specimen, that is, whether it was cold worked or annealed.

The pre-condition of the specimen had less of an affect on elongation and

reduction in area. The furnaces and their associated power supplies and

controllers were calibrated after much delay due to manufacture's errors

in the operating manuals (which were corrected by us in cooperation with

Leeds and Northrup Co.).

Stainless steel Type 310 was picked as a starting material for testing in the

constant strain rate machines because of its excellent resistance both to

corrosion and oxidation at high temperatures and because of its high

strength at elevated temperatures, thus making it a very desirable steel for

elevated temperature service. The 310 SS was first tested in an H 2 O/H 2

S gaseous mixture at 1000°F. Some of the early data indicated that SCC was occurring at slow strain rates in the H

2O/H 2 S environment. The same was

5
noted for 310 SS in a H 2

S environment in the absence of water vapor at the

slow strain rate. As we subsequently found, the 310 SS also failed by brittle

fracture with low ductility in a helium environment at 1000°F at slow strain

rates. It was then realized that what was earlier thought to be SCC was probably failure due to sensitization (precipitation of carbides or nitrides

in the grain boundaries).

A visit was made to the low BTU coal gasification pilot plant at Morgantown,

West Virginia to discuss the possibility of insertion of SCC specimens

in their plant.

Three talks have been given by Mr. Ugiansky concerning the constant strain rate

testing technique at the ASM/AIME Symposium on "Materials Behavior in Coal

Gasification Environments", at the NACE, Corrosion/76, Corrosion Research

Conference, and at the MFPG Symposium on "Prevention of Failures in Coal

Conversion Systems".

6
References

1. Heat Treating of Stainless Steel, Source Book on Stainless Steels, ASM,

pp. 162-173 (1976).

2. Uhlig, H. H. and White, R. A., Some Metallurgical Factors Affecting

Stress Corrosion Cracking of Austenitic Stainless Steels, Trans. ASM,

Vol . 52, pp. 830-846 (1960).

3. Troiano, A. R., The Role of Hydrogen and Other Interstitials in the

Mechanical Behavior of Metals, Trans. ASM, Vol. 52, pp. 54-80 (1960).

4. Hsiao, C. M. and Dulis, E. J., Effect of Interstitial Carbon Plus

Nitrogen and Precipitation Reactions on the Properties of Austenitic

Cr-Mn-C-N Steels, Trans. ASM, Vo. 52, pp. 855-877 (1960).

5. Shepherd, 0. C. and Schalliol, W. , The Effect of Environment on the

Stress-Rupture Properties of Metals at Elevated Temperatures, Symposium

on Corrosion of Materials at Elevated Temperatures, ASTM, Special

Tech. Publ. No. 108, pp. 34-41 (Apr. 1951).

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Fig.

4.

Stress-strain

curve

for

310

SS

in

vacuum

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line,

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Hg)

ata

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rate

of

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13

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310

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the

effect

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strain

rate

when

tested

in

dry

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(passed

through

an

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purifier)

at

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000°F
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Fig.

7.

Stress-strain

curve

for

310

SS

in

dry

helium

(passed

through

inert

gas

purifier)

ata

constant

strain

rate

of

8.43 •

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at

700°
700

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15

Stress-strai n

curves

for

310

SS

showing

the

effect

of

temperature

when

tested

in

dry

helium

(passed

through

inert

gas

purifier)

at

constant

strain

rate

of

8.43•

lO"' 7

S"l.
Fig. 9A. Photomicrograph of 310 SS as received. Etched for 5 seconds at

5 volts in 10% oxalic acid. 500 X

Fig. 9B. Photomicrograph of 310 SS after testing at 1000°F at a constant

strain rate of 8.43 • 10" 7 S~‘. Etched for 5 seconds at 5 volts

in 10% oxalic acid. 500 X

16
Fig. 1 OA . Photomicrograph of 310 SS as received. Etched for 5 seconds at

5 volts in 101 oxalic acid. 500 X

Fig. 10B. Photomicrograph of 310 SS after

strain rate of 8.43 • 10" 7 S' . in 10% oxalic acid. 500 X

testing at 700°F at a constant

Etched for 5 seconds at 5 volts

17
b) Erosive Wear (A. W. Ruff, L. K. Ives, J. P. Young, 312.03,

April - June 1976)

Progress : At this time tests conducted in the flame - erosion apparatus

have produced a fairly complete set of data covering the erosion of 310 ss in the velocity range 10 - 70 m/s using 100 mesh (150 ym) SiC abrasive

materials. These experiments were restricted to normal (90°) incidence

and were carried out both at 975°G in combustion gas atmospheres and at

25°C in air. A paper entitled "Erosion in 310 Stainless Steel at 975°C

in Combustion Gas Atmospheres" describing most of these results was prepared and has been accepted for presentation at the November 1976

ASME meeting in New York City and for publication in the J. of ASME.

A second paper describing some of the work was also presented at the

April MFPG meeting in Columbus, Ohio and is attached to this report.

Together the papers summarize the significant results that have been described

in previous quarterly reports with some significant additional information

obtained in the current quarter. A summary of the erosion data on 310 ss

is given in Fig. 1 where the log of erosion (sample weight loss in mg/

weight of abrasive in g) is plotted against the log of the particle

velocity. A similar graph was presented in the last quarterly report

(January - March 76). However, there an additional curve at 975°C

was included showing the erosion vs velocity dependence for SiC abrasive

material that contained impurities which resulted in the adherence of

abrasive particles to the sample surface at low velocities. The present

graph includes only data obtained with non-adhering SiC abrasive material.

In addition, redetermination of the characteristic particle velocity calibration curves of the erosion appartus has led to some changes with respect

to the earlier figure.

An important aspect of the elevated temperature studies has concerned an investigation of the morphology and composition of the eroded-corroded

surfaces. These investigations were carried out in a scanning electron

microscope equipped with an energy dispersive x-ray analysis system. The

specimens were subsequently sectioned normal to the eroded surfaces to

obtain information on the thickness, structure and in depth composition of

the oxide scale. A sample exposed at 975°C with a particle velocity of

14 m/s is shown in Fig. 2.

A brief summary of results - most of which are contained in the above

mentioned papers -is as follows. Considerable variations of oxygen - propane

gas ratio with associated changes in combustion gas composition at the

specimen location have little effect on erosion rate. At 975°C the erosion

rate is approximately proportioned to the velocity raised to the 1.8 power

while at 25°C in air the velocity exponent is approximately 2.5. At 15 m/s

the erosion rate is about an order of magnitude greater at 975°C than at

25°C. However, as a result of the different velocity exponents, the erosion

rates appear to converge at higher velocities. Over the velocity range

studied, 10 - 70 m/s at 975°C with a flux of 0.04 g/cm^ a relatively thick

oxide layer exists on the eroded surfaces. At the low velocity end of the

range, erosion was confined entirely to the oxide layer while at the higher

velocities particles were able to penetrate the oxide layer causing direct

damage to the underlying metal. The morphology of particle indentations

suggested a ductile flow of the oxide layer at 975°C.

18
A new series of tests was begun with the purpose ot measuring the angular

dependence of erosion at 975°C. However, in preparation to conducting

these tests some modifications were made in the flame-erosion appartus.

The most significant concerned the use' of tungsten-carbide (WC) nozzles

in place of the rapidly wearing stainless steel nozzles. It has been

necessary to replace the stainless steel nozzles after each test

400 g of abrasive) at low velocities and after 200 g of abrasive had

passed through the nozzle at high velocities 40 m/s at 25°C and ^ 70 m/s at 975°C) . Although precise data have not yet been obtained,

the life of the WC nozzles is substantially greater than that of the

stainless steel nozzles. The slightly smaller inside diameter (1.54 mm

vs 1.59 mm) of the WC nozzles has required the redetermination of the

particle velocity vs air flow curves.

Some preliminary measurements of angular dependence have been carried out at

a particle velocity of 34 m/s. Initially, simple wedge shaped samples

were used. However, it was found that the geometry led to a substantial

temperature gradient over the sample face. As much as 70°C was obtained

over a distance of 1.3 cm, the length of the exposed face, at a wedge

angle of 20°. Temperature measurements were conducted with an optical

pyrometer focussed on the sample surface. A thermal shield of cylindrical

shape was constructed to surround the specimen but this led to a reduction

of only about 20% in the temperature gradient. A more substantial improvement was obtained by employing samples of the geometry shown in Fig. 3.

A strip of 310 ss, 0.6 mm thick was carefully fitted over the top of the

specimen and clamped in place during exposure so that only the intended

specimen face was exposed to erosion. With this design at an average

temperature of 975°C, a difference of 25°C was measured between the top

and bottom of the 1.3 cm long face.

A few 975°C tests were conducted at attack angles of 20°, 30°, and 40°. In

all cases the erosion rate was signficantly less than at normal (90°)

impingement. However, the erosion rate at 20° was larger than at 30° and

40°. This suggests the existance of an intermediate minimum in erosion

between 20° and 90° followed by a low angle maximum between 0

° and 20 °.

Tests using the Roberts Abrader have been conducted for several purposes.

A new tungsten carbide nozzle (0.5 mm inside diameter) was placed into

operation and its performance was evaluated. A test was conducted at 25 C

using 304 ss material with CO 2 propellent and 50 ym A^O^, at flow rates

of 5 to 18 grams /minute. The relative weight loss and the relative cavity

depth formed at each flow rate was found to be reasonably independent of

flow rate as shown below.

19
Erosive Wear

Abrasive Flow

(g/min.)

Wt. Loss/Wt. Abr.

(IO- 3

) Penetration/Wt . Abr.

(ym/g)

5 0.17 6.6

8 0.19 7.1

12 0.17 6.7

15 0.19 8.3

18 0.19 8.6

This indicates that abrasive flow rates in this equipment is not critical

over the usual range employed.

Initial tests of the erosion resistance of thick chromium plate on steel

gave substantially larger wear resistances than those obtained for the

stainless steels and alloys containing nickel we have studied. A series

of tests was conducted at various temperatures, angles of incidence

and particle velocities. The abrasive used was 50 pm AI2 O 3 (90° incidence

for velocity tests) with CO 2 propellent gas for 2 min. exposure. Chromium

plate proved to have the largest erosion resistance in all the tests.

Some of the data are shown below along with those of comparative alloys.

Specimen

Particle

Velocity

(m/s)

Particle Velocity

Dependence of Relative Weight Loss

(10“3) at Indicated

Temperatures

25°C 500°C 1000°C

Angle

(deg)

Angle of Attack Dependence of Relative Weight

Loss ( 10“ 3

) a t Indicated

Temperatures and 31 m/s

25°C 500°C

Cr Plate 11 0.07 0.06 0.62 20 0.15 0.01

21 .11 .08 .61 30 .13 .02

31 .15 .11 .46 45 .15 .08

42 .14 .20 .73 70 .14

52 .16 .18 — 90 .17 .11

310 ss 11 0.08 0.09 --— 20 0.47

21 .15 .20 — 30 — .28

31 .18 .25 — 45 .30

42 .21 .33 — 70 — .34

52 .29 .55 — 90 — .14

304 ss 11 0.10 0.10 — 20 — 0.33

21 .19 .21 — 30 .35

31 .24 .30 — 45 — .26

42 .25 .38 — 70 — .23

52 .33 .29 90 — .26

The angle of attack dependence of erosion at 500°C is shown in Fig. 4

for several alloy in comparison to the chromium plate on steel. The

low angle peak loss is not found for the Cr plate, however, a gradual

increase in erosion with increasing angle is seen. 20
Three different types of chromium were compared for erosion resistance

in the Jet Abrader at 25 and 500°C. Rolled, ductile, ultra-pure

chromium and a chromium-low iron alloy were compared with thick chromium

plate on steel. The abrasive used was '50 ym AI 2 O 3 at 31 m/s velocity,

nozzle at 45 and 1 cm from the specimen. Average loss of three tests

for each type and temperature are shown below.

Unit Temp ( °C) Cr Plate Pure Cr Cr-Fe

Wt. Loss , , . Wt. Abr.

<mg/g)

25 0.13 0.14 0.26

500 0.05 0.11 0.21

Penetration . . . Ht. Abr. (lWg)25 3.8 8.9 6.6

500 4.9 7.5 5.5

Hardness (KHN^^ 25 985 230 587

The exact composition of the Cr-Fe alloy is not yet known but appears

to be a few percent of iron. The pure chromium and the Cr-Fe alloy were prepared by the Bureau of Mines several years ago.

The experimental Haynes 8077 alloy (16Cr, 4A1 + Ni + Y

2 O

3

) has been

tested in the Roberts unit at 500°C. AI2 O 3 (50ym) abrasive was used

at 31 m/s velocity with the nozzle at 45° angle and 1 cm from the

specimen. Test time was 2 minutes. The propellent gas was 003 . The

results were compared with those of several alloys previously tested

as shown below using the same conditions.

Unit Alloy Type

250MS 304 310 316 446 601 671 800 8077

Mean Wt. Loss (mg) 9.9 2.3 2.1 2.2 2.1 2.0 1.6 2.8 4.0

Mean Penetration (ym) 167 130 138 157 114 150 137 187 171

Wt. Loss , ,

. Wt. Abr.

(mg/g)

0.95 0.23 0.20 0.21 0.20 0.19 0.15 0.27 0.38

Penetration , , N Wt. Abr.

(ym/g)

16.1 12.5 13.3 15.1 11.0 14.4 13.2 18.0 16.4

The wear resistance of the 8077 alloy appears to be somewhat lower then

that of most of the other alloys at 500°C under these conditions.

21
Development continues on the electronic control unit to remove abrasive

flux variations during erosion test exposure. This feedback control

device senses the exit gas pressure at the nozzle and uses any variation

to adjust the vibrator drive amplitude on the abrasive chamber. The unit

is in operation at this time and adjustable parameters are being optimized.

Plans : The effect of sulfur contamination will be studied. Characterization

of scale deposits will be conducted. Measurements on the alloys of interest

will continue. The electronic flow control unit will be further developed.

22
Fig. 1 Velocity dependence of erosion at 975°C and 25°C.

Fig. 2 Cross-section of surface after exposure to

erosion at 975°C to a particle velocity of 14 m/s.

23
Weight

Loss

(nig)

/Wt.

Abrasive

(g)

Fig. 3 Erosion test specimen design for angular

dependence measurements

Fig. 4 Angle dependence of erosion at 500°C for indicated

alloys and chromium plate on steel

24
Proceeding* the Mechanical FalE.uA.eA Prevention Group

ColumbuA, Ohio [April 1976)

PARTICLE EROSION MEASUREMENTS ON METALS AT ELEVATED TEMPERATURES

L. K. Ives, J. P. Young, and A. W. Ruff

National Bureau of Standards

Washington, D.C. 20234

Abstract: Measurements of the erosive wear of metals resulting from

impact by abrasive particles under reactive environmental conditions

are being conducted. Specimens of several stainless steel and nickel

alloys have been exposed to a gas stream containing abrasive particles.

The particle flux, angle of attack, particle velocity, particle size,

surface temperature and other significant variables were controlled.

Different types of abrasives were used including A1 ? 0^, Si0

2

, SiC and

pumice. Two different erosion test apparatus have been developed and

used to obtain particle velocities from 10 to 90 m/s and temperatures

up to 1000°C. Different gaseous atmospheres have been established

during the erosion exposures. Weight loss measurements and penetration

depth measurements were conducted on the specimens after exposure.

Multiple alloy specimen packages were used in some experiments to obtain

relative erosion data. Erosion rates were determined in the temperature

range o£ 25 to 1000° C. Erosion depended on particle velocity over the

range v to V

-5

, increased with particle size, increased with particle

hardness , and was maximum for an angle of attack around 20° . The

effects of temperature and environment vary in a complicated way among

the different alloys.

Key words: Abrasive particles, Erosion, Erosive wear, Metals,

Oxidation, Wear.

Introduction : Studies of metal wear by erosion under various conditions

have been conducted [1-6] at room temperature and in a few cases at

elevated temperatures [7,8]. Industrial applications, for example in

the petrochemical industry and in new industries such as coal gasification [9], require data on metal erosion under conditions of high temperature and reactive environments. Much of that information still needs

to be acquired. Erosion of metals at room temperature has been found

to depend on particle velocity in the range v^ to v^ [1,2]. The angular

dependence of erosion exhibits a peak at about 20° angle of attack [1,4]

and the effects of different particle sizes have been investigated

[2,5]. Erosion can either increase or decrease with increasing temperature [7,8] depending on the metal and the atmosphere. The mechanisms

of erosion have been studied using single particle impact techniques

[10]. The present paper describes recent erosion measurements conducted

on metals at room and elevated temperatures up to 1000°C in different

gaseous atmospheres.

25
Experimental: Two different erosion testing apparatus have been used

in the present work. One system was based on an instrument developed

previously at this laboratory [11]. That instrument had been used

principally to determine the durability of coatings.* Considerable

attention was paid during its development toward achieving reproducibility of operation. Modifications introduced in the current program

have been designed to extend the operating capabilities to elevated

temperatures while maintaining that reproducibility. A schematic

diagram of the system is shown in Fig. 1. In order to obtain data at

elevated temperatures, the metal specimens were heated directly by

passage of a dc current through them during erosion testing. Current

values up to 200 A at 10 V were used to reach an operating temperature

of 500°C and above. The metal specimen size was typically 3 cm long x 1 cm wide x 0.5 mm thick. Battery clips (gold plated) were used to

achieve electrical contact at each end of the specimen during heating.

A thermocouple pressed in contact with the underside of the specimen

was used to determine the test temperature. The propellent gas-particle

mixture was also heated. After injection of abrasive particles, the

mixture passed through an eight foot length of 1.6 mm inside diameter

stainless steel tubing that was heated by the passage of up to 60 A ac current along its length. The gas -abrasive particle mixture was discharged through a right angle tungsten carbide nozzle, 0.5 mm inside

diameter. The exit gas temperature was adjusted so that the specimen

temperature did not change during the test although it may have been as low as 300°C at high flow rates when the specimen was maintained at

500°C. The propellent gas used in the tests reported here was 99.9%

C0o . The exposure time during each test was 2 minutes. Table 1 lists some of the characteristics of this test apparatus.

Table 1 Range of Operating Parameters for Roberts Erosion Unit

Gas: CC^jC^CO, mixtures

Particle sizes: 5 to 50 ym (0.5 mm nozzle)

5 to 150 ym (1.25 mm nozzle)

Gas Pressure: 20-70 psig (0.14 - 0.49 MPa)

Particle velocity: 10-60-m/s

Solids concentration: 0.5 to 3 g/1

? Particle flux: 0.5 to 3 g/cm -s Specimen temperature: ambient to 500°C+

Roberts Jet Abrader, Kameras Instruments. This instrument is identified only for the purpose of adequately specifying the experimental

procedure used.

26
The second erosion system has been described previously [12] and

permitted test temperatures of 1000°C to be achieved. Briefly, a propane -oxygen- air burner was used to direct a stream of hot combustion

gases containing entrained abrasive particles down a 41cm tube to impact on the specimen. The stream was approximately 2 cm in diameter. By

changing the oxygen and propane flows, the composition of the gases

impinging on the specimen could be varied. When oxygen was supplied in excess of the amount needed for complete combustion, an atmosphere com- posed of N

2

, O

2

, C0 ? and h^O was produced. When excess propane was

introduced, CO, H? and hydrocarbons were obtained in addition to the

above gases. Test times were usually 60 minutes. Table 2 describes

the operating features of this unit.

Table 2 Range of Operating Parameters for High Temperature Erosion Unit

Gas: Combustion gases; air

Particles: A^O^, SiC

Particle sizes: 50-200 ym

Particle velocity: 10-90 m/s

Solids concentration: 0.005 to 0.02 g/1

Particle flux: 0.04 g/cm^-s

Specimen temperature: ambient to 1000°C

Table 3

Parameters of Abrasive Particles Used in Erosion Tests

Type Size Hardness

(ym) (HK@25°C)

SiC 150 2500

A1 2°3 5 2100

Al

2 °3 50 2100

Sl0

2 5-50 820

F-Pumice 5-100 560

4F-Pumice 5-50 560

abrasives used in the tests reported here are listed in Table

Room temperature hardness values are listed. The A^O^ abrasives are metallographic polishing grade material having a narrow size distribution. A collection of the 5ym Al^O^ particles is shown in Fig. 2. The

particles passed only once through the system and were then discarded.

The particle velocity was measured using the rotating disk, time-of27
flight method [13]. Velocity measurements were conducted at the test

temperatures.

Weight loss values were determined by direct weighing after exposure

and cleaning. Microscopic inspection of the surface indicated whether

particle adhesion to the specimen was significant. If so, attempts

were made to mechanically remove the adherent particles by light brushing. Since scale formation occurred at elevated temperatures, it was difficult to remove all adherent abrasive and yet not affect surface

scale. Dimensional measurements of erosion crater depth and of specimen

thickness loss were made using microscopic or micrometer methods with a precision of the order of ±1 to 10 ym depending on surface roughness.

Examination of the eroded surfaces was conducted using optical and

scanning electron microscope (SEM) methods.

Results

Studies at 25°C and 500°C: Examples of isolated particle impacts are shown in Fig. 3 after a low flux exposure of an SAE 52100 steel specimen. The rim of each erosion crater consists of exposed material that

has been deformed plastically and whidh is susceptible to removal on subsequent particle impacts. That material is also more susceptible

to attack in a chemically reactive environment. Measurements have been

conducted at room temperature using the modified Roberts unit with

different abrasive particle velocities. Relative weight loss and penetration measurements on 304 stainless steel are shown in Fig. 4 for 90°

(normal) incidence, using two sizes of A1

? 0, particles. The erosion ! loss depends on the velocity raised to about the 2.8 power in these

tests. In general, we find that erosion is proportional to velocity in

the range of v

z,i> to v^ at room temperature. A greater erosion loss is seen here for the larger particles, 50ym A1

? 0,. The comparison between

weight loss and penetration distance is usually complicated by such

factors as erosion beam geometry changes when the abrasive type and size

are changed. It is preferable to use the same erosion measure when

comparing different metals.

The relative weight loss as a function of angle of attack a, the angle

between the specimen surface and the incident particle stream, has been

determined for 304 stainless steel at both 25°C and 500°C. The particle

flow and velocity were maintained constant as the specimens were exposed

at different angles to the particle stream. At 25°C a peak erosion loss occurs at about 20° followed by a gradual decrease as the angle is

increased up to 90°. The angle dependence of erosion at 500°C is different. The low angle peak occurs at about 10°, followed by a rapid

decrease and then little change occurs in the range 40-70°. The loss

appears to increase again as 90° impingement is approached. At 500°C

oxidation is more rapid than at room temperature, hence, the incident

abrasive particles experience a different surface condition. The

increased relative loss at 90° may be connected with the known erosion

characteristics of brittle materials where the peak loss occurs at

28
normal incidence, lhe oxide fiim or scale probably behaves in a more

brittle manner than the ductile underlying metal.

A group of eight different alloys have been examined at 25°C and 500°C

using this apparatus. Results on the erosion loss of the alloys using

two different sizes of Al-O, particles are shown in Table 4. The particle velocity was 30 m/s at 45° angle of attack using CC^ propellent.

Table 4

Erosion Measurements on Alloys Using Two Sizes of

A1 ? 0

3

Particles at 30 m/s and 45° Incidence

Wt. Loss(mg)/Wt.Abr. (g) Penetration (ym)/Wt.Abr. (g)

25°C 500°C 25°C 500°C

Alloy

Type 5 ym 50 ym 5 ym 50 ym 5' ym 50 ym 5 ym 50 ym

18Ni-250MS 0.08 0.12 0.04 0.95 10 9 0.5 16

AISI 446 0.05 0.05 0.01 0.20 10 6 0.6 11

AISI 304 0.05 0.47 0.02 0.26 7 26 0.2 12

AISI 316 0.04 0.10 0.01 0.21 7 9 0.8 13

AISI 310 0.03 0.09 0.02 0.20 5 9 0.8 13

Incoloy 800 0.06 0.32 0.01 0.26 8 20 0.4 18

Inconel 601 0.06 0.30 <0.01 0.19 8 20 0.4 14

Inconel 671 0.02 0.11 0.01 0.15 4 9 0.4 13

Both relative weight loss and penetration are shown. It is clear that

erosion increases significantly with particle size at both temperatures

and for all alloys. This size effect is particularly large for the

500°C tests where a decreased erosion loss is seen for all alloys for

the 5 ym particles (relative to 25°C) . Apparently the developing oxide

scale on the surface offers significant protection but only for the

small particle size. The data at 500°C using 50 ym size particles

varies among the alloys in comparison with the 25°C results. At 500°C

the alloys having larger chromium content in general show the least

erosion loss (although exceptions are present) . Oxide thicknesses

measured on specimen cross sections after testing at 500°C for 2 minutes

ranged from about 2ym down to less than 0.5 ym, depending on the alloy.

More data needs to be acquired at other exposure conditions of velocity,

particle size, and environment before definitive conclusions can be

reached concerning relative material erosion rates.

A comparison between the results obtained with different abrasives is

shown in Table 5. The data refer to 304 stainless steel in air at

29
Table 5 Erosion Measurements on 304 Stainless Steel Using Different

Abrasives at 30 m/s and 45° Incidence

Wt. Loss

Wt. Abr.

(mg)/

_(g)

Penetration (ym)/

Wt. Abr. (g)

Abrasive

(Size)

25°C 500°C 25°C 500°C

A1 2°3 0.5 0.2 7 0.2

(5 ym)

A1 2°3

(50 ym)

0.47 0.26 26 12

Pumice 4F

(5-50 ym)

0.21 0.04

1

29 2.7

Pumice F

(5-100 ym)

0.10 0.03 8 1.7

Silica 0.50 0.21 45 12

(5-50 ym)

30 m/s particle velocity at two temperatures, 25°C and 500°C. The increased erosion loss for 50 ym Al^O-, relative to 5 ym Al^O, particles

is apparent. The two grades of pumice show somewhat reduced losses

relative to 50 ym A^O, that are probably associated with the reduced

hardness value. However, a detailed study of these and other data shows

that factors other than abrasive hardness are significant. These

include "cutting efficiency" at elevated temperatures, tendency to fracture, hardness at temperature, etc. Such details can only by understood

as a result of comprehensive studies of the various abrasives. Another

relevant abrasive characteristic is the tendency for certain abrasives

to sinter to the eroded surface and hence offer protection.

Studies of the effect of surface hardness of the metal on erosion resistance are underway. The results of one series of tests are shown in

Fig. 5 using SiC^ abrasive and type 18Ni-250MS steel hardened to different levels. Relatively little effect on erosion rate is seen as the

metal hardness varies; that observation has been reported previously by

others [14].

Studies at 1000°C : Experiments conducted in the second erosion system

have used either 150 ym size SiC abrasive or 50 ym Al^O, particles.

Two different types of experiments were conducted. In one, multiple

alloy specimen packages were used in order to obtain data on relative

erosion for the alloys. In the other, a single alloy specimen was used

for each test. Figure 6 shows an example of a multiple alloy specimen

package after exposure in the test unit at 1000°C for 60 minutes using

30
150 ym SiC particles at a velocity of 50 m/s. Excess oxygen conditions

were maintained during the test. Each specimen in the package was 0.5

mm thick and was fabricated to have two surfaces exposed to the particle

stream at angles of 90° and 45°. Fiducial marks were placed on each

specimen surface in order to measure the dimensional loss after testing

One other edge, not exposed to the abrasive particles but exposed to the

heated atmosphere, provided information on the oxidation process. While

this specimen design provides an opportunity for obtaining relative data

between the alloy samples in the same atmosphere, there can be interfer- ences between neighboring specimens in the package, particularly if the

erosion rates differ greatly. In Fig. 6 the specimens in Inconel 601,

671 and 310 stainless steel exhibit the least erosion. On cooling to

room temperature the oxide scale spalled from the 304 stainless steel

specimen. Little or no spalling occurred on the other alloys.

Erosion data obtained from two tests using the multiple specimen package

under different environmental conditions are shown in Table 6. <

Table 6

Erosion Measurements on Alloys at 1020°C in Combustion

Gases at Approximately 55 m/s Using 150 ym SiC Particles

Equivalent Wt. Loss (mg)/Wt. Abrasive (g)

Excess Oxygen^ Excess Propane^

Alloy

Type

90° 45° 90° 45°

AISI 446 0.41 0.35 0.73 1.1

AISI 304 0.59 0.37 1.6 1.5

AISI 316 — — 1.2 1.3

AISI 310 0.32 0.28 0.58 0.64

Incoloy 800 0.42 0.32 0.56 0.64

Inconel 601 0.29 0.24 0.23 0.49

Inconel 671 0.26 0.28 0.15 0.25

1. Typical Analysis: 0

2 =13%, CO

2=10%, N

2 =64%, h 0= 13 *

2. Typical Analysis: h

2=8%, com, H

2 o=m, C0

2=5%

N

2=52%, 0ther=6% •

Examination of the results in the excess oxygen atmosphere indicate a small difference between the 45° and 90° surface erosion losses. However, under excess propane combustion conditions a greater difference

develops between the erosion rates at these two angles. The increased

loss at 45° associated with the ductile behavior is clear under excess propane conditions. The difference in erosion rate among the different

metals is also more apparent under those conditions. The most resistant

31
alloys appear to be those containing larger chromium concentrations.

That same characteristic appears to apply in more strongly oxidizing atmospheres. An example of the surface conditions that develop at 1000°C

after erosion is shown in Fig. 7a for 310 stainless steel exposed under

excess oxygen conditions. Much roughening due to impact damage is evident on the surface and some abrasive particles (or fragments) are seen imbedded in the specimen. A cross section through this specimen is shown

in Fig. 7b and indicates the oxide film present.

Summary

1. Two sets of test apparatus have been developed and applied to erosion

measurements on metals in the temperature range of 25°C to 1000°C. The

atmosphere, particle size and velocity, specimen orientation and other

significant variables can be controlled.

2. Significant effects on the erosion rates of the metals studied are attributed to the oxide film or scale that forms on the surface. The

chromium content of the alloy is one important parameter.

3. Abrasive hardness is a significant factor in erosion rate although

other abrasive characteristics may be equally important. Hardness

for a given alloy may not be an important parameter.

4. The maximum rate of erosion occurs at angles of 10° to 20° at temperatures up to 500° C, suggesting that cutting processes are a prominent

metal removal mechanism.

Acknowledgements

This study was partially supported by the Energy Research and Development

Administration/Fossil Energy and their support is gratefully acknowledged. Many useful discussions concerning these studies were held with

5. Wiederhom, D. E. Roberts, E. Passaglia, and others at NBS to whom

the authors are indebted.

32
REFERENCES

1. Neilson, J. H. and Gilchrist, A., "Erosion by a Stream of Solid

Particles," Wear, Vol. 11, 1968, pp. 111-122.

2. Goodwin, J. E. , Sage, W. , and Tilly, G. P., "Study of Erosion by

Solid Particles," Proc. Inst. Mech. Engrs. (London), Vol. 184, 1969,

pp. 279-292.

3. Kleis, I. R. , "Problems der Bestimmung des Straheverscheisscs bei

Metallen," Wear, Vol., 13, 199 (1969).

4. Sheldon, G. L., "Similarities and Differences in the Erosion Behavior

of Materials," Trans. ASME, Vol. 92_, 1970, pp. 619-626.

5. Finnie, I., "Some Observations on the Erosion of Ductile Metals,"

Wear, Vol. 19, 1972, p. 81-90.

6. Tilly, G. P., "A Two Stage Mechanism of Ductile Erosion," Wear,

Vol. 23, 1973, pp. 87-96.

7. Tilly, G. P., "Erosion Caused by Airborne Particles," Wear, Vol. 14,

1969, pp. 63-79.

8. Smeltzer, C. E., Gulden, M. E. , and Compton, W. A., "Mechanism of

Metal Removal by Impacting Dust Particles," Trans. ASME, Vol. 92,

1970, pp. 639-654.

9. "Workshop on Materials Problems and Research Opportunities in Coal

Conversion: Conclusions and Recommendations," Staehle, R. W.

,

Director, Ohio State University, April 16, 1974.

10. Hutchings, I. M. and Winter, R. E. , "Particle Erosion of Ductile

Metals: A Mechanism of Material Removal," Wear, Vol. 2_7, 1974,

pp. 121-128.

11. Roberts, A. G., "Improved NBS Abrasive Jet Method for Measuring

Abrasion Resistance of Coatings," ASTM Bulletin No. 244 , 1960,

pp. 48-51.

12. Wiederhom, S. M. and Roberts, D. E., "A Technique to Investigate

High Temperature Erosion of Refractories," Bull. Am. Ceram, Soc.

,

Vol. 55, 1976, pp. 185-189.

13. Ruff, A. W. and Ives L. K. , "Measurement of Solid Particle Velocity

in Erosive Wear," Wear, Vol. 315, 1975, pp. 195-199.

14. Finnie, I., Wolak, J. , and Kabil, Y. , "Erosion of Metals by Solid

Particles," Journal of Materials, Vol. _2, No. 3. 1967, pp. 682-700.

33
VIBRATION

DRIVER ABRASIVE

CHAMBER

SPECIMEN

MOUNT

FLOW AND

Fig. 1 Schematic representation of

the modified Roberts erosion appartus.

Fig. 2 Examples of 5 ym sized abrasive particles.

34
Fig. 3 Erosion impact craters on steel surface.

Fig. 4 Velocity dependence of erosion

loss at 25°C for two particle sizes.

35

Penetration

(/im)

Wt

Abrasive

(g)
CP

<D>

took_ -Q<

Fig. 5 Variation of erosion loss for different

hardness levels in type 18Ni-250MS steel.

Fig. 6 Multiple alloy specimen package after erosion test

at 1000°C. Two exposed surfaces 45° and 90°, are shown.

36
Fig. 7

(a) Surface of 310 stainless steel after erosion test at 1010°C.

(b) Cross section of specimen showing oxide film and erosion craters.

II H

'' 37
B. Ceramics

(a) Deformation and Fracture (E. R. Fuller, Jr., C. R. Robbins and J. M. Bukowski,

313.05)

Progress : High-pressure steam, one of the reactants used in the gasification of

coal, has been found to interact with the bonding phases of several castable

refractories used as liners in coal gasifiers. In some instances, these interactions disrupted the bonding structure, thus, adversely affecting the structural

integrity of the castable. To examine these effects, measurements of the flexural

strength of these materials have been performed at room temperature subsequent

to hydrothermal exposure. Mineral phase transformations resulting from the

hydrothermal treatment was monitored by x-ray diffraction analysis. High-pressure

steam was found to have a dramatic effect on a high-purity, calcium-aluminatebonded castable and on high-purity, calcium aluminate cements. The flexural

strength of these materials decreased sharply after hydrothermal exposures at

temperatures above approximately 350 C. This degradation of mechanical properties

appeared to be related to the formation of a high-pressure, calcium-aluminate

hydration phase, C^A^H^ in the cement matrix tha^ bonds the alumina aggregate,

and to the dehydration of other bonding phases. In contrast to this behavior,

a calcined flint clay castable was found to be mechanically unaffected by

hydrothermal treatment. Additional bonding phases in the cement matrix appeared

to maintain the integrity of this castable for the hydrothermal conditions which

were examined.

The room temperature bend strength subsequent to hydrothermal treatment for these

two commercial refractories is plotted in figures 1 and 2 versus the hydrothermal

treatment temperature. The nature of this effect has been examined further in

the past quarter. Polished sections of these commercial castables were examined

subsequent to hydrothermal exposure using both optical and scanning electron

microscopes. Chemical compositions of various mineral phases were determined with

energy dispersive x-ray analysis. These results were correlated with the mineral

phases that have been identified by x-ray diffraction analysis. See Table 1.

These transformations are the result of chemical reactions between the highpressure steam and the ceramic components of the castable refractories. After

firing and prior to hydrothermal treatment, the high-purity, calcium-aluminate

castable refractory contained only alumina and calcium-aluminate phases. Hydrothermal

treatment results in the formation of several hydrated alumina and calcium-aluminate

phases. As the hydrothermal treatment temperature increases these phases undergo

several dehydration steps (as listed in Table 1) . Until the formation of the

C^A^H^ phase, these hydration phases are ones commonly observed during mixing and

curing of these castables. The strength degradation appears to be associated

with the formation of this high-pressure hydration phase of calcium aluminate and

Cement chemistry notation used throughout this section is:

A = A1

2 0

3

; C = CaO; S = Si0

2

? and H = H

2

0

38
to the dehydration of boehmite (AH) , which was formed in the refractory at lower

temperatures. With the dehydration of C AH

& and the formation of C A H , the

boehmite apparently acts as a predominant bonding phase in the matrix of the

refractory, providing substantial strength to the matrix. However, when this

phase decomposes, the bonding matrix loses its structural integrity which appears

to cause the decrease in the after-exposure- strength. Examination of polished

sections with a scanning electron microscope revealed the crystallization of the

C^A^H^ phase in the bonding matrix, thus disrupting the bonding.

In contrast to this behavior, the calcium-aluminate cement with calcined flint

clay aggregate did not react with steam to form predominantly this hydrated

calcium-aluminate phase. Instead high-pressure steam appears to accelerate the

reaction between the calcium-aluminate phases which are initially present, and

the fine calcined flint clay aggregate to form a calcium alumino-silicate

,

anorthite (CAS

2

> , that maintains the strength of the bonding matrix. The

anorthite, a feldspar mineral commonly found in igneous rocks, forms a strong

continuous structure that penetrates throughout the refractory matrix and into

some of the larger aggregates firmly holding these coarse aggregates in place.

The high strength of this calcined flint clay refractory is attributed to this

structure of the matrix.

These environmental effects on the bonding matrix have been further verified by

examination of sized fractions of these commercial refractories (+150 mesh and

-150 mesh) . The -150 mesh fraction of each refractory were cast into specimens

and their flexural strength determined after exposure to hydrothermal environments.

See Figure 3. Again we see that the high-purity, calcium-aluminate refractory

exhibits a strength degradation; whereas the calcined flint clay castable maintains

its integrity, confirming the previously determined results. In addition to

examining the cement matrix, effect of hydrothermal environment on the alumina

aggregate have been investigated. The room temperature bend strength of three

densities of sintered alumina were examined subsequent to hydrothermal treatment.

See Figure 4. With increasing hydrothermal treatment temperature each type of

sintered alumina exhibited a strength loss that appears to be related to the

density, or porosity. However, no mineral phase changes were detected from the

x-ray diffraction analysis. It is speculated that steam attacks the grain

boundaries to form grain-boundary hydration phases, thus weakening these sintered

alumina specimens.

These studies demonstrate the importance of chemical reactions and phase transitions

in castable refractories intended for use as liners in coal gasification plants.

Reactions that form weak phases seriously degrade the mechanical properties of the

liner, and significantly decrease its expected lifetime in service. Thus far,

the importance of water as a reactive species has been demonstrated. However,

since other gases in coal gasification environments are likely to cause similar

degradation of strength, it is important to identify chemical reactions that can

39
occur between these gases and the refractories. Toward this objective three

pressure vessel systems are presently being developed: two for exposure testing

with subsequent mechanical property characterization; and one for mechanical

properties testing under coal gasification conditions. The specifications and

properties of these pressure vessels are listed in Table 2. System 2 for exposure

testing is a commercial system which has been ordered. The projected delivery

date is the latter part of July 1976, The other system for exposure testing

(system 1) is a modified in-house pressure vessel. It is presently operational

and is being used with CO^ and/or steam in experiments described below. However,

further minor modifications are necessary for the introduction of other gaseous

environments and for the venting of obnoxious gases. A gas manifold and venting

system are being developed for this purpose. A major consideration in the

construction of the manifold and venting system is the ease of adaptability to

the other two pressure vessel systems. Preliminary design of pressure vessel

system 3 for in situ mechanical properties testing has been completed. Basically

this system is a two pressure vessel system. The inner pressure vessel is a hot-walled pressure vessel containing the coal gasification environments. This

vessel is contained in a room temperature pressure vessel which will be pressurized

to the same pressure so as to support the inner vessel. Long-lead-time items for

the construction of the outer vessel have been ordered.

To provide some quality control over the refractories which are tested, three

laboratory refractory compositions are being prepared for these studies. These

compositions - two calcium-aluminate-bonded and one phosphate-bonded - are typical

of the spectrum of refractories expected for use as liners in coal gasification

systems. The two calcium-aluminate-bonded refractory compositions will be

similar to the two commercial refractories previously studied. One will contain

tabular alumina aggregate with a high-purity calcium aluminate cement; and the

other will contain calcined kaolin aggregate with the same calcium aluminate

cement. The third refractory composition, a phosphate-bonded refractory, will

contain a tabular alumina aggregate and will be bonded by phosphoric acid. The

components to form these refractory compositions have been obtained and the

refractory aggregates graded into various size fractions. The calcined kaolin

refractory composition has been prepared from calcined kaolin aggregate and

calcium-aluminate cement in the following weight percents: coarse aggregate - 65% calcined kaolin (+ 325 mesh) graded into various size fractions; fine

aggregate - 10% calcined kaolin (- 325 mesh) ; and 25% of a calcium-aluminate

cement. The resulting refractory has a chemical composition of approximately

56% alumina, 37% silica, 4.5% lime and 0.8% iron oxide. Refractory specimens,

prepared from this mixture and fired according to standard practices, have a control flexural strength of approximately 1,300 + 350 psi. Investigation

of the mechanical properties after exposure to CO^ and /steam environments

has begun for this composition. Experiments performed thus far were run at

610 C and 1000 psig total pressure. With a pure atmosphere, 65 hour

exposure resulted in little change in strength; the after-exposure bend strength was 1,200 + 220 psi. However, when the environment consisted of 50% CO^ and

50% steam the strength approximately doubled. This has been verified in both

a 20 hour exposure test and a 165 hour exposure test which had after-exposure

40
bend strengths of 2,170 + 380 psi and 2,180 + 440 psi, respectively. X-ray

analysis of the chemical phase is still in progress at this time. Preliminary

results suggest that this increase in strength is related to an enhanced

crystallization of anorthite (CAS 9

) and CaCO^ (CT) .

Plans : Mechanical properties testing after Exposure to environments of CO^ and

C0o

/ steam at elevated temperature and pressure will be continued. Design and

construction of pressure vessel barricades, gas manifold, and venting system for

obnoxious gases will also be continued. These systems will be applicable to

all three of the pressure vessel systems. The design of the internal pressure

vessel and the loading system for the in situ system (system 3) will be completed,

components ordered, and construction begun on this part of the system.

41
Table 1. Predominant Phases, as Determined by X-Ray Diffraction Analysis,

Subsequent to Hydrothermal Treatment. Phases Listed in Apparent Order of

Abundance.

i

i

Temperature

High-purity alumina

castable

Room Temperature a-Alumina (a-A)

CA , CA

2

110°C (230°F

)

a-A, C

3

AHg,

Gibbsite (AH^)

210°C (41 0°F

)

a-A, Boehmite (AH),

C

3

AHg, C

4

A

3

H

3

(trace)

310°C (590°F

)

350°C (662°F) ) 375°C (707°F)

J

a-A, C

4 A

3 H3

41 0°C (770°F) i 510°C (950°F)

j

a-A, C 4

A

3 H3

Calcined flint clay

castable

Mullite (AgSg),

Cristobal ite

,

a-Alumina (a-A), ca

2

, Anorthite (CAS^),

CA, Amorphous Phase

Cristobal ite, A^S^ct-A,

Gibbsite (AH^),

C^AHg, CAS 2>

Cristobal ite, A

3

$

2

, Boehmii

(AH), a-A, C

3

AH

6

, CAS

2

,

C

4

A

3

H

3

A

3

S

2

, Cristobal ite,

a-A

,

AH, C

4

A

3 H3

, CAS

2

, C

3

AH

SOI

A

3

S 2> Cristobal ite, CAS

2

, a-A, Tridymite

42
Table 2. Specifications and Properties of Pressure Vessels

For Exposure Test For in situ Test

System 1 System 2

/ System 3

max. temperature 600°C 960°C 1 01 o°c

pressure limit at

temperature 5,100 psi 3,000 psi 1 ,000 psi

gauges (max. working

pressure)

0-600 psi

(450 psi

)

0-1500 psi

(1125 psi)

0-1500 psi

(1000 psi

)

0-2000 psi

(1500 psi)

0-3000 psi

(2250 psi)

Dimensions

volume 1 ,000 ml 260 ml

internal diameter 2.5" 1 .3" 4

"

internal length 13.13" 11.85"

length of hot zone - 6" ~ 4" ~ 6

"

# of vessels 1 2 1

source of steam

pressure

closed system;

generated from

h

2 o

water

pump

pressurize

water

system

projected

installation

May 76 Aug. 76 Jan. 77

projected

operation

June 76 Sept. 76 April 77

43
;

HIGH PURITY, CALCIUM -ALU MIN ATE CASTABLE

50 230 410 590 770 950

HYDROTHERMAL TREATMENT TEMPERATURE

Figure 1. Room Temperature Bend Strength versus Hydrothermal Treatment Temperature for

a High-Purity, Calcium-Aluminate Castable

CALCINED FLINT CLAY CASTABLE

(°FI 50 230 410 590 770 950 1130

Figure 2. Room Temperature Bend Strength versus Hydrothermal Treatment Temperature for

a Calcined Flint Clay Castable.

44
TEMPERATURE ( F)

Figure 3.

Figure 4

33

31

29

27

. 25

23

21

|

19

l

17

i i5

50 230 410 590 770 950 1 130

i

.

13

11 .

9 & - Calcined Flint Clay

» - Tabular Alumina

CONTROL 10 110 210 310 410 TEMPERATURE (°C)

510 610

Room Temperature Bend Strength versus Hydrothermal Treatment Temperature for

-150 Mesh Fractions of a High-Purity, Alumina Castable and of a CalcinedFlint Clay Castable.

so

o

a.

SINTERED ALUMINA

<°F)

100

- 35 000

- 30,000

- 25 000

- 20,000

- 15,000 CONTROL 10 110 210 310 410 510 HYDROTHERMAL TREATMENT TEMPERATURE (°C)

"O

tf>

. Room Temperature Bend Strength versus Hydrothermal Treatment Temperature for

Three' Densities of Sintered Alumina.

45
b. Erosive Wear (S. M. Wiederhorn and D. E. Roberts, 313.05)

Progress : Castable refractories will be used in coal gasification applications

where resistance to both chemical attack and erosion is required. Applications

will include the linings of major components of gasifier plants such as reaction

vessels, hot transfer lines, and cyclone separators. To assure adequate performance in these applications and to accumulate information that can be used

by engineers for design purposes, this program is being conducted to characterize

the erosion resistance of castable refractories. An important goal of this project

is to determine the type of castable refractories that will be both erosion

resistant and sufficiently inexpensive to satisfy the economic requirements of

coal gasifiers.

With these objectives in mind, this project was initiated to characterize the

erosion behavior of castable refractories. In the initial stages of the program

an apparatus was developed which could be used to study the high temperature

erosion behavior of refractories, in environments that simulate coal gasification

environments. The apparatus is capable of temperatures that range from 25 C to

1000 C and of accelerating particles to velocities of from 10 to 100 m/s. This

range of temperatures and particle velocities cover the range expected in

gasification plants. Although the apparatus was designed to operate at one atmosphere pressure, the environmental composition can be regulated by controlling

the input gas composition. An analysis of the gases at the specimen site indicated

that the gas composition was similar to that of the Bureau of Mines' gasifier at

Morgantown, Pa.

Early studies using this apparatus indicated a rather strong dependence of erosion

rate on particle velocity. The erosion rate could be expressed as a power function

of the particle velocity, with a power exponent of approximately 3. This value is

similar to that reported for other materials. Erosion occurred primarily by wear

of the cement phase of the castable refractories in such a manner that channels were

often observed to form in the surface of the refractory. As the cement phase was worn away, aggregate grains would fall from the refractory surface resulting in

general wear. By increasing the strength of the cement phase (by reducing the water

to cement ratio for example) channel formation, and thus wear could be suppressed;

conversely, wear could be accelerated by adverse chemical reactions (for example

high pressure steam which has been shown to weaken some refractories)

The erosion resistance of a high purity alumina castable refractory was found to

increase slightly as the temperature was increased from 25 C to 1000 C. This

increase in erosion resistance was accompanied by a change in the erosion mechanism.

At 25 C erosion occurred by a brittle fracture process during which small particles

of the refractory were chipped from the refractory surface by the abrasive particles.

At 1000 C the erosion was controlled by plastic deformation. Evidence for plastic

deformation was obtained from optical microscopy and scanning electron microscopy

studies, and from investigations of the erosion rate as a function of particle

Impingement angle. At high temperatures great numbers of erosion particles were

observed to become imbedded in the surface of the refractory indicating that the

temperatures have softened the refractory to permit a degree of plastic flow.

46
Further evidence for high temperature plastic flow was obtained from studies of

erosion as a function of particle impingement angle. At 25 C the dependence of

erosion rate on angle of impingement. Figure la, is typical of that found for

brittle ceramic materials such as porcelain: the erosion rate is maximal at

90 impingement and decreases slowly as the angle of impingement decreases. At

1000 C, the erosion behavior, Figure lb, is similar to that reported for metals:

the erosion rate is minimal at 90 and increases as the angle of impingement

decreases passing through a maximum at approximately 20° before decreasing to zero. As will be seen in the discussion given below, this high temperature

behavior is modified by the development of surface channels.

During the past year the results discussed above were extended by investigating

the effects of hydrothermal environments on the erosion of castable refractories.

Two types of refractories were used in the investigation. One was a high purity

aluminum oxide refractory consisting of tabular alumina bonded with a calcium

aluminate cement. The other was a calcined flint clay aggregate also bonded with

calcium aluminate cement. Cubes of these refractories, 1/2 inch on a side, were hydro thermally treated by sealing them in a fixed volume (20 cc) Morrey bomb that

contained approximately 3 cc's of water. The bomb was placed in a furnace which

had been pre-heated to the temperature of reaction. After 20 hours of exposure

the bomb was removed and permitted to cool to room temperature. Specimens

were then removed from the bomb and dried in air. The erosion studies were

conducted at room temperature to determine if the erosion properties of these

materials had been changed by the hydrothermal treatment. Each specimen was

eroded by 100 grams of 100 mesh silicon carbide particles, at velocities that

ranged from 15 to 50 m/s. The erosion rate was determined from weight loss

measurements after exposure.

This study demonstrated a susceptibility of the high alumina castable refractory

to chemical attack under the conditions of test. Below 200 C exposure, the

erosion resistance of the high alumina castable was essentially unaffected by the

hydrothermal treatment (Figure 2a) . At approximately 300 C the erosion rate of

this castable refractory increased by a factor of 10, indicating a significant

change in its erosion resistance. By contrast the erosion resistance of the

calcined flint clay castable was unaffected by the hydrothermal treatment over

the temperature range 25 C to 500 C (Figure 2b) . A comparison was made between

the erosion resistance and the strength of castable refractories that received

similar hydrothermal treatments. Over the same temperature range, the strength

of the high alumina castable refractory decreased precipitously at 300 C. Whereas

the strength of the calcined flint clay castable refractory was unaffected by

hydrothermal treatments.

In order to determine the cause of the change in strength and erosion resistance,

an X-ray analysis was done of the castable refractories after the hydrothermal

treatments. The strength degradation and enhancement of the erosion rate appear to be

associated with the formation of a high pressure hydration phase of calcium

aluminate (ACaO.lA^O^.Sl^O) and the dehydration of boehmite which forms in the

47
refractory temperatures between 110°C and 210°C. With the dehydration of

3Ca0.Al

2 0y6H9 0 above 210°C and the formation of 4CaO. 3A1

? 0

3

.3H

2 0, the

boehmite apparently acts as the predominant bonding phase in the matrix

and provides substantial strength to the refractory. However when this

phase decomposes, the bonding matrix loses its structural integrity

resulting in the observed decrease in both the strength and erosion

resistance of the refractory.

In contrast the above behavior, the calcined flint clay refractory did not

react with steam to form a low strength hydrated calcium aluminate phase.

Instead the calcium aluminate that was initially present reacted with the

mullite aggregate to form calcium alumino-silicate phases that maintained

their strength. One of the calcium alumino-silicate phases was anorthite,

a feldspar type of mineral commonly found in igneous rocks. Although highly

porous, the calcium alumino-silicate phases formed a strong structure that

penetrated throughout the refractory firmly holding the mullite aggregate

in place. It is this matrix that accounted for the high strength and erosion

resistance of the calcined flint clay refractory.

This study demonstrated the importance of chemical reactions and phase

transitions to mechanical properties. Reactions that form weak phases will

seriously degrade the mechanical properties of linings and significantly

decrease their expected lifetime in service. In this study the importance of

water as a reactive species was demonstrated. However since other gases in

coal gasification environments are also likely to cause degradation of strength

and erosion resistance, it is important to identify the chemical reactions that occur between these gases and the refractories. Therefore one objective of

future work will be to identify other chemical reactions that are detrimental

to the mechanical integrity of the refractories.

In more recent work, the effect of long term exposure on erosion rate was also

investigated. This study was conducted to determine the effect of surface

topology on the erosion rate. As erosion occurs in refractories, the cement

matrix that holds the aggregate particles together is slowly worn away. The

result of removing matrix material is to leave open channels within the castable

refractories. In earlier phases of this project small amounts of abrasive were

used and as a result wear in the matrix material was often not fully developed.

Therefore it was felt essential to determine erosion rates on castable refractories

under conditions in which fully developed channels formed in the castable refractory

so as to evaluate the erosion properties of these refractories after long term use.

The study of long term wear was conducted by exposing castable refractories

to an erosive stream at relatively high velocities, 70 to 90 m/s. The high

velocities were used so that channeling would develop in a relatively short

48
time. The study was conducted both at 25°C and at 1000°C . In order to determine

if the erosion rate changed as the surface topology evolved, the experiment was conducted using increments of 25 grams of abrasive for the first 100 grams of

exposure, 50 gram increments for the next 200 grams of exposure and 100 gram

increments for the final 300 grams of exposure. Each specimen was exposed to a total of 600 grams of abrasive. In order to determine if the erosion rate

depended on the incident angle of particle impingement, the castable refractories

were exposed at angles ranging from 15° to 90 . Specimens for these angular

studies were pre-cut so as to expose the surface at the correct angle.

In both the high and low temperature studies the erosion rate was found to initially

decrease as the amount of abrasive used was increased and then to approach a steady

state (Figure 3 and 4^. The change in the initial erosion rate was not as great

for the^studies at 25 C as for those at 1000°C. For impingement angles of less

than 45 , there was little change in the erosion rate at 25°C as the amount of

abrasive was increased. By contrast at 1000°C the erosion rate for the first

25 gram increment of exposure was considerably greater than that observed in

subsequent exposures. This change in the rate of erosion can be attributed to

the development of a steady state surface structure. For the room temperature

studies the surface remained relatively flat during exposure. By contrast the

surfaces at 1000 C had deep channels around the abrasive particles.

The results of the 25°C study differed from an earlier one in which deep channels

were observed to form around the aggregate particles. It is felt that this

difference arises from the method of making the refractories, which had changed

from the earlier method. Refractories are now made using a cement mixer and

small amounts of water, whereas the earlier refractories were hand mixed

using larger amounts of water. The change in the erosion rate of the

material suggests that an increase in the water-to-cement ratio reduces the

erosion resistance of the cement, permitting channels to form. This finding is

in agreement with earlier findings in the literature. The results of the current

study also suggest that if the water-to-cement ratio can be reduced sufficiently,

then the cement matrix can be made nearly as erosion resistant as the aggregate

particles, resulting in a greater resistance to erosion.

At high temperatures, however, the fact that the cement matrix becomes plastic,

reduces its effective erosion resistance, and results in the formation of deep

channels. It is interesting that these channels are formed regardless of the

direction of impingement. The structure of the channels reflect the angle of

impingement. Thus at 90 exposure the channels go straight down into the

refractory while at low angles of exposure, the channels lie at an angle to the

surface. The presence of erosion resistant aggregate particles have the effect

of modifying the appearance of the refractory surface and of protecting the

surface from erosion.

49
The effect of angle of impingement on the erosion rate is shown in Figure 5 for

studies at 25°C and 1000°C. The erosion rate changes only slightly with the

impingement angle. At room temperature the erosion rate decreases by a factor

of approximately 20% as the angle impingement is changed from 90 Q

to 15°. At

1000 C the erosion rate is relativgly constant between 90 and 30 , and then

drops by a factor of about 2 at 15 . The low temperature behavior (Figure 5a)

is similar to those reported earlier (Figure la) where only 100 grams of abrasive

were used, and where relatively low particle velocities were used (20 m/s). At

1000°C the behavior in Figure 5b differs considerably from that shown in Figure lb.

It is believed that the shape of the curve obtained in this figure reflects

the fact that the surface topology was different in the two studies. The

development of deep channels in the present study modifies the erosion behavior.

While the erosion still occurs by a ductile flow mechanism, the formation of

channels and steps on the surface effectively results in a 90 exposure angle

for all angles used. This is illustrated in Figure 6 where the erosion was shown

schematically in two dimensions. During the very initial stages of erosion the

surface is flat but as the cement matrix erodes away pockets are formed between

the aggregate grains. Each of these pockets is aligned normal to the direction

of the incoming particles. Therefore the total effect of channel formation in

the refractory is to form a surface that lies at 90 to the impingement particles.

As a consequence this behavior the erosion rate now becomes insensitive to the

angle of impingement. This behavior can have significant practical importance

if it holds for other castable refractories, since erosion measurements will not

have to be made over a range of impingement angles to characterize the wear of

the refractory. Instead a convenient angle can be selected to measure the erosion

rate and to completely characterize the erosion resistance as a function of

impingement angle. This result further implies that in estimating the rate of

erosion for castable refractories in service one can assume the rate of erosion

to be dependent of impingement angle.

Plans : During the next year further erosion studies will be conducted to

characterize the erosion rate of castable refractories. Target materials will

include a phosphate bonded tabular alumina refractory and a calcined Kaolin refractory.

The effect of temperature on the erosion rate will be investigated by exposing

refractories to 100 grit silicon carbide particles at temperatures of 25 and 1000 C.

Particle velocities will range from 10 to 100 m/s and erosion measurements will be

conducted using 90 impact angles. The effect of angle on the erosion rate will

also be tested for a selected range of angles. Study will not be as extensive as the one just completed unless some difference in behavior is noted.

Since erosion in gasifiers will be due to particles that are much softer than

silicon carbide, studies will also be conducted to evaluate the effect of particle

hardness on erosion rate. Dolamite, silica sand, tabular alumina and ground-up

coal slag will be used as erosion particles. Studies will be conducted at

temperatures of 25 and 1000 C using an impingement angle of 90 . In order to

determine the effect of particle size on erosion rate these materials will be sized

to form erosion mixes having particle diameters of approximately 50, 150 and 500

micrometers.

50
In addition to the above studies the hot erosion equipment developed earlier

in this project will be suggested as a standard ASTM technique for characterizing

the erosive wear of ceramic refractories. The technique will be described at a meeting of ASTM Committee C-8, and if interest is expressed in the technique

arrangements will be made to have similar sets of equipment constructed at other

laboratories. Round Robin Tests will then be conducted to determine the accuracy

of the equipment and its applicability to erosion tests. The adaptation of this

technique will provide a uniform standard by which the erosion resistance of

castables can be measured.

1

51
Erosion

Halt,

micregrams/flrsm

of

Abrasive

Figure 1. Initial stages of wear. Dependence of erosion rate on the

angle of impingement (a) 25°C; and (b) 1000°C. The error

bars represent one standard deviation from the mean.

52
EROSION

RATE

(milligrams/gram

of

abrasive)

HYDROTHERMAL TREATMENT TEMPERATURE

(Degrees Celsius)

HYDROTHERMAL TREATMENT TEMPERATURE

(Degrees Celsius)

Figure 2. Effect of hydrothermal treatment on the erosive wear of

castable refractories. Tests conducted at room temperature

after hydrothermal treatment.

53
A

Figure 3. Low temperature erosive wear as a function of the amount of

abrasive used.

54
Ei~o%ion

(wj/gm

o~f

a

hhci

$>/ J

ft)

A

too 7.00 ^oo Ijoo 500 &oo

Tct*l A bt-qsiJe

Figure 4. High temperature erosive wear as a function of the amount

of abrasive used.

55
Xrr\ p

i n<^c. rf\ef)T Anyle. TmpiHgC’tncnt Ang/e

Figure 5. Erosion Rate as a function of impingement angle. Long term

erosion with full channel development. High purity aluminum

oxide castable refractory bonded with calcium aluminate. The error bars represent one standard deviation from the mean.

56
Erosion Direction

Figure 6. Schematic diagram of channel development. Left-hand figure shows

castable refractory prior to erosion. Right-hand figure shows

the early stages of wear in which the cement phase is preferentially

worn, leaving channels on the surface. Since the bottoms of these

channels are aligned perpendicular to the erosion direction, erosion

occurs at the same rate as if the original surface were aligned

perpendicular to the direction of erosion.

k

57
2 .

a.

Chemical Degradation of Ceramics

Gas Phase Reactions

(W.S. Brower,

313.03)

J.L. Waring and C.A. Harding,

Some work in the CaQ-Al-O^-H^O system has been previously reported in

the literature. The most outstanding is the work of Peppier and Wells

and Majumdar and Roy. For the purposes of studying the solid state reactions

encountered with steam in a coal gasifier the solubility work of Peppier

and Wells was reinterpreted with the view of integrating the data with that

of Majumdar and Roy. This was accomplished by constructing the diagram

given in Figure 1 from the data in Table 1. From Figure 1 it can be seen

that the various existence regions can be generally characterized. From

this preliminary data it was found that the work of Peppier and Wells

and Majumdar and Roy can be integrated to predict behavior at 1000°C.

Therefore it was necessary only to verify that portion of the system which

could directly effect the operation of a coal gasifier at 1000 psi of

steam at 1000°C.

A constitution diagram giving the results of this study (Figure 1) has

been constructed from the data given in Table 1. From Figure 1 it can be

seen that the first hydrate formed 3Ca0:Al 2 0- :6H

20 is unstable and dissociates above room temperature to form CafOH) and 4Ca0:3Al

20~ :3H20.

Preformed of 1000°C at 1000 psi steam. The 1000°C/1000 psi data was obtained

with the TZM apparatus. The maximum and minimum dissociation temperature

for the 12Ca0:7Al

20

2 hydrogarnet have not been established. However the

minimum temperature of formation is probably in excess of 400°C and the

compound should be stable to the reported melting point of about 1392°C.

The hydrate 4CaO: 3A1 9

0> : SH^O has a minimum temperature of stability below

200°C and a maximum temperature of stability below 600°C in 1000 psi steam.

The Ca0:Al

2 0„ and Ca0:2Al2 0_ were found to exist at temperatures in excess

of 1000°C ana to hydrate Below about 400°C. The compounds CaOtA^O^,

CaO:2Al

2 0^ and Ca0:6Al2 0^ experimentally appear to have the same minimum

temperature of stability since at these relatively low temperatures the

reactions proceed very slowly and it has not been possible to delineate

these boundaries, therefore the region of uncertainty is shaded on the

diagram. At a temperature below about 400°C-450°C these phases dissociate

into boehmite and 4CaO : SAl^O^ : 3^0 . The phase transition between boehmite

and Al

2 0g in the pressure range of 1000 psi is currently under investigation.

In conclusion at 1000 psi of steam a refractory cement consisting of a mixture of Ca0:2Al

20», Ca0:Al2 0~, and A12 0~ grog might be expected to react

slowly at about 450°C. The SCaOrA^O-^l^O originally formed in the cement

will transform rapidly to form 4Ca0:Al202:3H

20, which will then slowly react

to form first 12Ca0:7Al

20^ plus CaOsA^O^ ana then eventually CaOrbA^O^ plus

ai

2°

3

.A similar investigation was conducted in CO under 860 psi. Although

insufficient data was collected to construct a pnase diagram, the data given

in Table 2 indicates that the reactions which occur in a C0

2 environment are similar to those which occur in steam. There are several exceptions. The

4CaQ:3Al

2 0

2

:3H

2O at 500°C in steam is stable and under similar conditions

58
in C0 o decomposes to CaCO^ and Al^O^. Several experiments indicate that

other^Ca0-Al 9 0_ phases may also be stable with respect to the CaCO -A1„0

system under these conditions.

Two specimens consisting of CA-25 and 4Ca0:3Al 0 :3H 0 were heated

in a methane environment at 800°C for two weeks at 1000 psi (Table IV)

.

The x-ray powder pattern of CA-25 (see Table III) showed Ca0:2Al

20„,

Ca0:Al

2 0

2> a-Al^^ and an unknown phase (major line ^24 deg 20). The

unknown phase does not appear to be calcium formate. The 4Ca0:3Al

2 0 :3H 0

contains 12Ca0:7Al

20

2 and CaOrAl^O^. These are the same reaction products

as those occurring under 1000 psi steam at the same composition and temperature.

Four compositions were heated in CO (Table V) at 1000 psi and 700°C

for the x-ray pattern of Ca0:Al 0

3

showed a mixture of Ca0:Al o 0

3 and

Ca0:2Al

20

3 and the 4Ca0 : 3A1

2 0~ :3H

2 0 showed the same decomposition products

found in the C0

2 and CaCO^ . The CA-25 x-ray pattern showed mixtures of

CaCO^, ot-Al

20

3

, Ca0:2Al

20

3 and CaOtA^Q^ and the Ca0:2Al

2 0

3 showed no reaction.

Conclusions

It is obvious that insufficient experiments have been conducted to

make any definitive conclusions. However the following observations

should be noted.

1. At 1000 psi steam pressure, the phase boundaries are essentially

those which can be predicted from a combination of the published data of

Peppier and Wells and Majumdar and Roy.

2. The cementations phases 3:1:6 and 4:3:3 are not stable at any

higher temperatures in 1000 psi steam than they are at atmospheric pressure.

3. The compound 4:3:3 decomposes to metastable phases in dry atmospheres

such as C0

2 or CO but directly to stable phases in steam.

4. Many of these phases in the CaO :A1

20

3 :H

20 system are not stable

in C0

2 (860 psi) with respect to the system CaCO^-A^O^.

In order to determine the possible deleterious effects of the combined

gaseous atmosphere present in a coal gasifier considerably more detailed

study must be made of this system.

b. Slag-Refractory Interactions

A synthetic slag was formulated by averaging the ash compositions of

twelve Pittsburgh coals. (Bur. Mines Bull. 567) with some simplification

arrived at the following composition.

SiO,

Al

2°

3

Fe 2°3

CaO

MgO

k

2 o Na

20

49.0% by wt.

26.0

19.0

3.5

0.7

1.2

0.6 59
A large batch of the above composition was prepared and the viscosity

determined over the temperature range 1620 °C to the crystallization temperature of 1470 °C; Figure 3.

A series of magnesium aluminates were prepared in the form of right

cylinders 1 cm diameter by 2.5 cm long for immersion in molten slag. The

compositions are as follows:

Density

A^MgO^ Stoichiometric Spinel 3.45 g/cc

90.1 mol% Al

2Mg04 9.9 MgO

74.1 Al

2Mg0

4

25.9 MgO

18.6 Al

2Mg04

81.4 MgO

3.40

3.50

3.46

A cylinder of stoichiometric magnesium aluminate spinel was immersed in

the molten slag to a depth of 3/4 - 1 cm and rotated without reversal at

15 rpm for two hours at a temperature of 1550 °C. The sample was sectioned

and visually there was little or no penetration of Fe into the specimen,

dissolution was minimal and was estimated to be less than 0.20 mm or 0.1

mm/hr. The same sample examined by SEM did not show any Fe beyond the

frozen slag layer, indicating that the degradation process was primarily

by solution at the slag-ceramic interface. A sample of 90.1 mole % MgO.A^O,,-

9.9 mole % MgO was subjected to the same slag composition in the same manner for a period of three hours at a temperature of 1575 °C. The cross sectional area had a diffusion layer which could be visually observed as well as the original outline of the sample. The diffusion layer was approximately 1.3 mm thick corresponding to a diffusion rate of 0.13 cm/hr.

In a likely manner the end member composition 18.6 mole % MgO.Al^^-

81.4 mole % MgO was subjected to the slag environment under the same conditions of time, temperature and rotation. The diameter was reduced

to 0.778 cm with no obvious diffusion of Fe into the sample. The rate of

solution is approximately 0.36 mm/hr. The degradation of the stoichiometric spinel is primarily by solution with a very small diffusion layer.

With increasing amounts of MgO first diffusion predominates until the

solution rate becomes equal to or greater than the diffusion rate. The

obvious conclusion that can be drawn from this is that stoichiometric

spinel is more resistant to slag attack than spinel - MgO compositions.

From phase considerations one would expect that the stoichiometric spinel

would be more resistant than A^Q^ under the same conditions.

Plans : The reactions of molten slag with refractories has essentially been

discontinued. The effort will be directed toward studying the properties

of real and synthetic slags under simulated gasification environments to

establish a base of information useful for engineering design.

3. Consultation and Proposal Review

Progress : During the year, thirty proposals were reviewed for ERDA-FE.

Six of these proposals were reviewed in the fourth quarter.

60
to

LO CM

CO CM

;sd lunSS3Ud

61

TEMPERATURE
Temperature C

62
TEMPERATURE

(C)

63

Figure
Starting

Material

Temperature

Pressure

Time

X-ray

Diffraction

Analysis

Conclusion

°C

psi

days

V

ro + co c O O 0

(D CM ro 0 CM e co

•P r—

i

O -P t—

1

•H O

•H ft CM •H ft p CM e CM rH £ CM 0 <—

1

JS • • < X • • Po ft

(1) o o • • 0 O o X • • O cm (T5 o 0 CM cd 0 o

X! K u cd X! K u 0 0

ro u ro — c o

+ + t t o- 0

ro 1 co XJ + roO ro coo CO 0

O CM O ro O CM O cop ro cmh CM O C\H CM O O

i—

l

rH CM rH ft t—

1

CM0 CM < ro ft i—

1

ft CO < H X) r—

1

CM •• CM < CM •• CM ft ft

•• O • • •• o • • •• X

O 0 o 6 O (d O O 0 6

0 u (0 rd cd u cd 0 0 0

u M' u u U M- O O E u

+o

CM

ro -p

• • G

ro ,—.. •H O 0-. 0

CM s— + 4H

r-l CM 0

ft • • p 0

ro i—

1

•rl >

t. E O 0

O 0 X CM x: 0 o 0 K

U cd O co rH p X 0

+ -P + CO E

ro CO + coo CO O o CO O CM O co COM CM0 IM O CHH CM O o ••

i—1 -P 11 CM r—\ ft. i—

4

CN CMH

ft *H ft i—

1

ft to ft i—

1

rH

CM £ CM ft CM •• CM c ft 0

•• X. • . O • • •• u O 0 o 6 O rd o 6 o 0

0 0 0 rd cd O 0 id 0 P O XI U u O M’ u O U -p

r- r- r- rH rH rH 1—

1

t—

I

rH 1

1

1—

1

o o o o o o o

o o o o o o o

o o o o o o o 1—

1

1—

1

rH rH rH I—

1

rH

o O o o O o o

o O o o o o o

ro r- r- in O' in

ro oro o ro

ro Oro O ro ro

CM CM O CM CM O O rH rH CM rH rH CM CM

ft ft t—

l

< ft rH rH

CM CM ft CM CM ft ft 6 6 O 6 6 O 6

0 0 0 0 0 0 0

u u u u u u u

64
Table II. Experimental Data for the System CaO-Al^-l^O at 1000 psi Steam Pressure.

Starting.^

Material—

Temp.

°C Pressure

psi

Time

hrs X-ray Diffraction

Analysis Conclusions

3:1:6 350 1000 720 4:3:3 + Ca(0H)

2 + hydrate 3:1:6 4:3:3 + hydrate

3:1:6 450 1000 720 4:3:3 + Ca(0H)_ + 3:1:6 + hydrate 3:1:6 -* 4:3:3 + hydrate

3:1 350 1000 720 3:1 + Ca(OH) + 12:7 3:1 -+ 12:7 + Ca(OH)

2 3:1 450 1000 720 3:1 + Ca(0H)

2 + 12:7 3:1 -* 12:7 + Ca(0H>

2 3:1 500 1000 168 3:1 + 12:7 + Ca(0H)

2 3:1 12:7 + Ca(0H>2 3:1 600 1000 168 12:7 + 3:1 + CaO 3:1 12:7 + CaO

3:1 700 1000 168 3:1 + 12:7 + CaO 3:1 12:7 + CaO

3:1 1000 945 120 12:7 + CaO 3:1 12:7 + CaO

4:3:3 350 1000 720 4:3:3 + boehmite Ca(0H)„ + amorphous A1

2 0» +

4:3:3 -* 4:3:3 + boehmite J

4:3:3 400 1000 86 4:3:3 + Ca(OH)

2 4:3:3 + amorphous Ca(OH) + amorphous Al-O. -*• 4:3:3 +

Ca(OH)

2

4:3:3 450 1000 720 4:3:3 4:3:3 + amorphous Ca(0H),

; + amorphous A1

203

-+ 4:3:3

4:3:3 500 1000 86 4:3:3 + Ca(OH)

2 4:3:3 + amorphous Ca(OH) + amorphous A1„0, 4:3:3 + cryst Ca(0H)

2

4:3:3 500 1000 168 4:3:3

4:3:3 + amorphous A1 0 +

amorphous Ca(0H)„ 4:3:3 +

Ca (OH) 4:3:3 600 1000 168 4:3:3 + 12:7 + CaO 4:3:3 + 12:7 + CaO 4:3:3 700 1000 168 (leaked) 12:7 + 1:1

1:1 300 1000 168 4:3:3 + 1:1 + boehmite 1:1 -* boehmite + 4:3:3

1:1 400 1000 240 1:1 1:1

1:1 450 1000 240 1:1 + 4:3:3 1:1 -> 4:3:3

1:1 500 1000 264 1:1 1:1 -*• 1:1

1:1 700 1000 168 1:1 + 1:2 1:1 + 1:1

1:1 1000 945 120 1:1 1:1 1:1

1:2 300 1000 168 1:2 + 4:3:3 + boehmite 1:2 * boehmite + 4:3:3 +

boehmite

1:2 400 1000 264 1:2 + boehmite + 4:3:3 1:2 -* boehmite + 4:3:3

1:2 450 1000 240 1:2 1:2 1:2 1:2 500 1000 264 1:2 1:2 -+ 1:2 1:2 700 1000 168 1:2 1:2 -*• 1:2

1:2 1000 945 120 1:1 1:1 - 1:2

1:6 350 1000 945 1:6 + 4:3:3 + Al^ 1:6 + 4:3:3 + boehmite + A1

203 1:6 450 1000 945 1:6 + A1

2 0

3 1:6 -+ 1:6 + A1

20

3

- 3:1:6 = 3CaO :A1

2 03

: 6H

2 0 1:2 = Ca0:2Al

2 0

3 3:1 = 3Ca0:Al

20

3 4:3:3 = ACa0:3Al

2 0

3 :3H

2 0 12:7 = 12Ca0:7Al

2 0

3 1:1 = Ca0:Al

20

3 1:6 = Ca0:6Al

2 0

3
Table

III.

Experimental

Data

for

Compositions

in

the

System

Calcium

Oxide

(Carbonate)

Aluminum

Oxide

in

860

psi

of

C0_. j

COdo

•H

CO3

i—I Odou

CDtn iH

COd

cx

uo

CO rH rH + + + i + f 1 + + 1 v© VO vO CO m CO CO

iH rH rH rH iH CO CO CO CO

CO CO CO CO CO CO sr St *t <r

co Ou

CO o+cn oCM J-H

<d+

CO oOCO u+CN

N

N « i

CM CM fv© o x)l q1 A d CM CO Q

42} 0 • © 0 c

rH

O CO

O 3i CO

d

d CO CO CM + cO

O

d + 0 CO

rH CJ 0 0 O

® E V

u

M CO CJ 0 3 cO 0 0 iH + « d 0 cO CM

co

d e0 so CO + 0 COu 3 0 V rH

CU 14

H

VV • • rs

d 0 + + cd + <3

Ctf <d + CO 0 • 4*S CM-- co

u CO CO l •H rs rs • A CM d rH T>l O + 0 0 a O 0 • e • <0- rH 3 < P 0 CM CM CJ f CM CM + + T + 5 CO A • cO d

to rH rH CO CO CO cO

O

u rH <jO

d

CO + + 0 0 0 •• *5 + + + + + d sH rH rH 0 0 CJ CO J2 rH rH rH rH CM rH CH CM 6 co c0 c0 •• fi X CO CO CO u u CJ 'O- 3 ?H rH rH rH rH rH rH rH rH c0 O B-5 CM iH O O 0 O O O O 0 O 0 0 0 0 0

rH 32

O vO vO vO CO CO CO CO c0 0

•Hd in

CU r-'. rH d CO vO

eg 0 &$ • • a CM f-H m m m O O O 0 0 O 0 0 0 0 0 vO m 1—1 O CM CM CM rH rH rH CO CO CO CO m m m m VO co

00 < 0 dd CO d CO c0 O iH • m d cO O m m m O O O 0 0 O 0 0 0 0 0 CO QO

co CJ 0 rs rs rs CO CO CO -a- MJ- <r <fr

m

m

mm CO CM

cu 42 cu 42d d

0) 0) S d

<u e

T)

u o d d

CO CO M-l d

"O 51d

<U 0 O X *H -H rl U (0 0 <U d

a. <u 51 C0 g d tu cuMua)

cu <u

5)

d 4

5-i

H cO 3 S a • ^ CJ C0 42

*H r—I

42 cO 'U £X*H <U CO d 42 d CU U

SOW

C

O cO CU d 0 3 c b -h

CO

d

dO

5) 4J

O

4-1 co cu

d 0 d -h

(30

w 43

d co 0

•H

d

O

<D 43 a

44d

cO

4-> CO

cO

<ud3d . CO

d rH Xi

d

d CO cu <U <u d u CX • •H 0 co 0 'd O •H 0 3 cu •H d •H d d <4-1 d X) CO HH <u d 0 CU 3 > •H O 42

CO

d

O 0 CD d CD

d 0 CU CO

co

d

d S

co 3 cO xt d

d

(U

d cO

d •H

d

d 0

•H (U CD *H 0 X) cx d 0

rH cu 0 •H <U

CO

d s 4-( CX

O CO

d CD

<u

d 4*i 'd cu H d 42 3 cO 4-i 42 O

CU

d

H io H ^ C d 42

d <U 3 42 0 d 1 45O

4-1 °rl CO 42

CU £

4-» cO

O

<D 4-1 d

do q>d3 d 4-> d co

cu

d

co cu

<u CX B. I4-» dd <u 3 42 O d

CU 42d

<u CM CM -a- vO •<r 00 ^r St -t 00 CM CM StO CMO

d or CU

d fv. <r CO 00 -a- 00 00 <r St 00 St CM On CMO CU <4-4 CD 42 1—i

CO CO rH CM CO iH rH CM 1—(

rs iH CU CD cu

O <U

H CX cQ

d

d

CD 42 <U

d cx

CX & CU CU <u -d <y d d (U x» d CD 3 •H <U rH d O 0

CO CD

O

O 0 0 0 OO 0 O

OO 0 0 OO

OO 15 00 CO 42

CD cx VO vO vO vO V© v© V© V© v© v© V© v© V© v© v© V© V©

dd

d

<U CO 00 00 00 00 00 00 00 00 00 00 00 00 00 00 00 00 42 -H <u •H d O CD 42 ts

en •H 3 d rH

CU $ O 0 O O 0 0 0 0 0 0 O 0 0 0 0 0 0

CD rHd

43

(X 00 CDCD

O 0 O 0 0 0 0 0 0 0 O 0 0 0 O 0 0 CO

O

d CU <U V© fs. rs in V© 00 m V© 1". 00 m V© rs 00 m r'- m •H cu d O

cO CUd

x)

<U

d

id

•H

CU

d

d

cu CO X>

•H CO

CU CO CO

CU co

CO 42

CO CX 42a cuCO

<u cu 42 42 H H

66
33

<ucn3

a)3

<u5

co3O

4-1 U3

33O

J-t CO 0 a U 0 CM o CM 32 3 3 33 CO 0

vO CO • • •H

4-1 •• 0 CO 4-1 3 CO CO CM 0 co co 3 O 0 0 rH CM CO O O 3

-3 CM CM < H 0 CM CM 33

CO r-H rH c CM

H

H So

C < • • CO

H < <d 43

00 • • • • 0 • • < CM 0 3 3 0 0 3 O 33

•H 3 3 u 3 0 O 0 5 u 0 CM U 3 3 3 rj o co co rH 0 U u 3

1—

0 H II 11 11 II 11 II 11 3 o H M3 co H CM \£> 3 • • 3

<u CO rH CM coHH

H

3; • • rH • • 3 H CO <3* <

o\

T3 | 67
Table

IV.

Experimental

Data

for

Compositions

in

the

System

Calcium

Oxide

(Carbonate)

Aluminum

Oxide

in

Methane.

CO oCM

rH t < o oCM CO

CO 33

O 3 CO -f O • < O")

p

•H COO

O 3 M

CO

O CM cu

O 33 3 3 CM—1 0 Li 4-1 I—1

rH <3 •3 c 00 o <3 4-> <U 4-4 33 3 3 CO •• PL 3 O •H o .. O 4J O •H Li o O CO c /-n 33 Li

cO o CU

C 4-1 S 3

CJ CM •H CO 3 O

'S’

H1 O Li

O

O

•H

C tH 4-1 o

CM *H CM

CM 3 3 3 33 M Li 3 3 <U 3 3 *H

4-1 3 4-1 rH

4J CO

O 3

CO CO <U 1 u Li

CO

O 33 3 3 3

•H CM 33 PL T3

CO rH cu cu 3 0 £

tH < 3 M 3 3 8

i-H I •H 0) 4-1 4-1 CL

CO

a CO

U £ 3 3 + o rH 3 3 3 < CO CM 3 CO Li 33 o o rH O 3 00 4-1 •H CM <3 CU 'W 4-1 o 1 o rH cu i 4-1 O m *h < O Li 43 4-1 3 3

4-1 4-t cO <U

U 3 Li rH O o C-> 5 <u 3 4-1 cm 3 <0 cO + PL 3 3 CM

CO

M

u CO <u co 3 3 •H

0) LH + o Li Li 3 T3

Od CM PV CM 3 <U

a 3

•H o oj rH PL CO

m tH

Q CMC <3 0) 4-1 CL 3

rH 3 i"" o a 3 Li tn <c o • • •H H 3 3 « CO CM 5 O 33 0 3 X M •• 5 C0 CL i ° 1 o c O 3 • 33 3 J3 Li X c0 3 CM Li CO g 4-1 0 U + rH 00 rH 0 3 n O 3 33 33 to 3

Li *H 4-1 rH 0 4J M rH iH 0 3 3 Cm Li 3

11

U 1 O 3 33

ft

O

w 3 4-1 CO 0 «U Li 3

<U vO v£> 33 3 3 33 0 Li i—1

rH CU (U 4-1 •3

O 4J •H 33 CO CO Li CD O Li 3 •H H CU 3 M O 3 £ 33 XJ

PL to 3 4-1 3

33 33 •HO 3

0)

O <U 3 3 H •H H H 3 33 3 o o 33 00 3 3 3 CL

CO *H o O S tf 3 > Li

CO 00 o o O •H 3 p 3 CL rH rH 3 3 tH 00 n rH M 3 O

PL, 3 3 3 3 3 3

•H *H 3 Li 3 3»S Li 3 > 3 3 3

111

u

p 33 3 • 4-4 33 O 3 PL

PL

O o 3 0 0 3 3 0 U O o 0 •H 3 TO 3 3 o 00 00 O 3 cm 3 3 H 00 4-i •H 3 4-1 3 3 3 44 33 3 3 • •H 3 tH 3 H 3 •H CD

4-4 Li 3 33 3 CM o Li 3 3 X) 4-1 CM 3 W 3 •H • 3 3 00

33 4-13 3 3 3 O 3

CO to

L 44 3 Li £ O X) • • 3 3 3 \ CO PL 0 3 4-13 3 -S

col o -H

a 3 3 3 3 3 CM & CM 3 3 3 33 Li 0 0

00 fH rH 33 4-1 -H CL 3 iH •rH 4-1 3 CO <3 4-1 CL O o 3

"H CO

H 3 3 3 0 3 3 O

4J Li

m • • rH 3 rH 33 3 CL CL 33

M 0) CM o <3 3 0.

H 41 3 C/3 3

(0 4-) 1 cO

4-1 CO

O co a o «l £>\ o|

68
Table

V.

Experimental

Data

for

Compositions

in

the

System

Calcium

Oxide

(Carbonate)

Aluminum

Oxide

in

Carbon

Monoxide. i

CO OCN

»H

<3

.. cn O o 3 CN

CJ+ s 3 • • c o o o CM 3

•H 33 3 CJ 3 CO S 3 .. o + (—

co3 o O co c CMC o o •H 3 CN o <3 + rH i co co

..

O s O CJ • • 3 3 o

CJ CJ 3 1 + CJ 1

CO •H CO

CO

O

X CJ CO

rH 3 o 3 CJ CN C + rH

<3 co <3 O CM 3 CN °L ..

Q

O rH

a

O

CO *H <3 3 3 3 J-) CM O CJ

rH

O .. 3 + 3 3 O CO

CO 3 CO 3 O 3 4H CJ 3 CN

pc! M-I + + CO rH CO

•H CO co o <3

O

Q o o CN 1 CN

CM CN rH 8 -H

X rH rH <3 + <3

CO <3 <3 CM CO' • 3 • • • • o o 1 O O O CJ 3 X cO 3 3 3 CJ a CJ CJ CJ +

0) o OOO 6 3 iH rH rH rH

•H 4= iH rH l—

rH

H

<u33 O

O o O

CO -HO o o O

CD CO

O o o o 3 a. tH rH rH rH 3

CH

a. o O o o 0 CJ o O O o

a) o r-- r-- r-~ H O

CM

33co• •cn

cOl o CO

CO CO CM

O

60 *h

O i—

CM 3 cO CM <3 i—

•H -H rH co <3 H 3 <3 • • CN

m 3 3 O • • CN

CO Hi 6 3 O 1 HI CO 3 CJ 3 <3

C/3 S3 CJ CJ cj oo 3 3 3 S .3

•H 3 3 3

HI 3 3

3. 3

HI

O y—

3 3 3 X 3 3 3 3

*H 3 3 u •H

O

•H 3 UH 3

4h •H <3 3 3 3 d 3 3 3 3 d d 3 o 3

HI 3 Q 3 3 3 3 3 X 3

33 3 a 3 3 3 0 3 3 3 3

•H 3 3 3 O & 3

rH 3 3 3 CO 3 X 3 3 60 3 3 3 e 3 3 •3 3 3 3 y 3 3 3 3 3 a 3 3 3 co 3 3 3 3 3 3 3 (X 3 a CD 3 3 3 3. a X 3

•H

H

d 3

.3 o 3 a • 0 O 3 • rd 3 X 3 CO g 3

60 rH o 3 O 3 X X

X 3 •H 3 rH

HI 3 rH •H O 3 3 4H 3 3 3 B O 3 a 3 c 3

co 0 3 3 3 X 3 3 3 3 3 -a U 3 3 o 3 3 3 3 3 o 3 & X 33

ex

X 3 3 x X •tH O y 3 3

•H rH rH 3 x 60 3 3 3 > 3 3 > 3

•H

O •H 3 3 3 •H 60

rH 3 3 3 3 3 3 3

•H •H 3 3 3 3 3 > 3 3 3 3 3 4= 3 X O d CX 3 O U 3 . 6 •H 3 d O 3 3 3

60 3 •H 3 3 3 3 3 X 3

•H 3 *H 3 H 3 3 3 3 X 3 d 3 d 3 3 X •H • 3 3 3 d 3 3

CO 3 3 3 3 S 3 3 d

r*

a

g 3 3 3

rH s 3 3 3 3 d 3 3 X 3 0 X 3 •H a 3 •H 3 CX U i—

d 3 3 g 3

rH 3 rH 4= 3 Cu c 3 <x H 3 3

x\

gocG3

<3

y|

69

dehydration

product.
4. Failure Analysis (J. H. Smith, 312.01)

A. Failure Avoidance Program

Progress : Additional reports of pilot plant operating discrepancies and

component failure have been received during this quarter. Approximately

250 items of information have been received from nine different pilot

plants and process development units. All of this information has been

evaluated, classified, and abstracted for inclusion in the Failure Avoidance Program data base.

As part of the information collection phase of this task, visits were made to the ERDA-Synthane pilot plant. Bigas pilot plant, and the Westinghouse combined-cycle pilot plant at Waltz Mills. The operating

experiences and component failures experienced to date were reviewed

with the plant operating personnel at these pilot plants. At present,

the Bigas Pilot Plant has not started operation, so no information was received from this source, but arrangements were made to receive available

information on component failures when operation of the plant is started.

The Synthane plant has had limited operation of some subsystems which

resulted in a few shutdowns due to component failures. Details of these

failures were reviewed and the information in the Failure Avoidance Program

data base were brought up-to-date. The Westinghouse-Waltz Mills pilot plant

has a very detailed system for documentation of operating discrepancies and

component failures. This system was reviewed in detail and arrangements

were made to study the information available from this source.

Additional information collection activities included attending the

American Society for Metals Conference entitled "Materials for Coal Conversion Systems Design" on April 26 and 27, 1976 in Pittsburgh, Pennsylvania.

In addition, activities of the Metals Properties Council Committee on Materials for Coal Conversion were participated in.

The evaluation, classification, and preparation of detailed abstracts has

been completed for all information received to date from the pilot plants.

An automated, computer based data system has been established and developed

to permit efficient storage and retrieval of all information included in

the data base. At present, all available extended abstracts and classification of information items has been put on this system. The use of this

system is being refined and training of personnel in its use is continuing.

The use of this system will permit rapid and efficient access to all failure

avoidance information and will permit failure mode analysis programs to be

conducted.

As part of the information dissemination phase of this task, a presentation

entitled "The ERDA-Fossil Energy Failure Avoidance Program" was presented

at the Mechanical Failures Prevention Group Meeting on "Prevention of

Failures in Coal Conversion Systems" in Columbus, Ohio. In addition, a

70
summary of results obtained to date was published in the April 1976

issues of the ERDA-Materials and Components Newletter.

A presentation was made to explain the use of the Failure Avoidance

Program at the annual GIDEP information meeting on May 27 in Alexandria,

Virginia. In addition, the possible use of the GIDEP system for future use in the Failure Avoidance Program was explored.

Plans : The information gathering phase of this task will continue as more information is received on a continuing basis from the pilot plants.

In addition, detailed follow-up of certain information already received

will be initiated. Additional visits to operating pilot plants will be

made to collect additional information.

The information evaluation and proce-s phase of this program will continue

and be enlarged. The operation of the automated data storage system will

be enlarged and refined for easier and more efficient use. The information

presently in the data base will be used to conduct failure mode analyses

to identify critical operating problem areas.

The information dissemination phase of this program will continue and

will be developed to include more direct contact with users. At present,

several inquiries a month are answered from plant operations and design

personnel. Plans have also been initiated to include summary results of

this program in the bimonthly ERDA-Materials and Components Newsletter.

Additional information dissemination will take place through the Fossil

Energy - Performances Assurance Systems that is currently under development. Close liaison with this program is being maintained.

71