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ORNL-2524
Metallurgy and Ceramics
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THE INFLUENCE OF COMPOSITION UPON

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THE 1500°F CREEP-RUPTURE STRENGTH AND

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MICROSTRUCTURE OF MOLYBDENUM-

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CHROMIUM=IRON~NICKEL BASE ALLOYS
T. K. Roche

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UNCLASSIFIED ORNL-2524

Contract No. W-7405-eng-26

METALLURGY DIVISION

THE INFLUENCE OF COMPOSITION UPON THE 1500°F CREEP-RUPTURE
STRENGTH AND MICROSTRUCTURE OF MOLYBDENUM-
CHROMIUM-IRON-NICKEL BASE ALLOYS

Thomas Kirby Roche

DATE ISSUED

JUN < 41958

Submitted as a Thesis to the Graduate Council of the University of Tennessee
in partial fulfillment of the requirements for the degree of Master of Science

OAK RIDGE NATIONAL LABORATORY
Operated by
UNION CARBIDE CORPORATION
for the
Atomic Energy Commission

UNCLASSTIFIED

¥ SYSTEMS LIBRARIES

[N

3 445k 0361261 N

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-if-

ACKNOWLEDGEMENT

The author is indebted to Dr. E. E. Stansbury for his advice
throughout the course of this investigation and for his cortributions
to the preparation of the final maruscript.

Special thanks are due H. Inouye and D. A. Douglas, Jr., of the
Oak Ridge National Iaboratory for their helpful comments.

Direct assistance in accumulating the experimental data for this
work was provided by the following ORNL personnel: G. E. Angel, melting
and casting; J. F. Newsome and W. R. Johnson, fabrication; C. K. Thomas,
creep-rupture testing: W. H. Farmer, metallography; W. R. Laing and
staff, chemical analyses. Their contributions are gratefully appreci-
ated.

The author also is indebted to Mrs. Freda Finn of the Metallurgy
Reports Office for her cooperation in typing this manuscript.

Finally, thanks are extended to Union Carbide Corporation for

its provision for employee educational assistance.

 
-iii-

TABLE OF CONTENTS
CHAPTER ' PAGE,

T. SUMMARY. & + ¢ « o v o o o o o o o o o o o v o e e m e v e 1

II. INTRODUCTION « « « o o v o o o o o « o o o o s o o v o a o L
IIT. OBJECTIVE:. « « + « o o o o o o o o o o o o o o e s v v v 10
TV. EXPERIMENTAL PROCEDURE « « « o o « o & o « o v o + « o o . 11

V. RESULTS AND DISCUSSION + + + o « o v ¢ v o o o s + « o « . 23

VI. CONCLUSIONS AND RECOMMENDATIONS. . +» . + « « « o 4 4 4 . . 89
LIST OF REFERENCES. + « « « v o v « & o ¢ + o o v s v s v o v v v+ 93
BIBLIOGRAPHY. « » « o « o o o « o o « o + o o v v o s o v o o u v O5

APPmDIX . * L] . . - - * - . & * » . [ ] . . . . - . . . . . - » * . . 97

 
CHAPTER I
SUMMARY

Results of an alloy development program at the Oak Ridge
National Iaboratory have shown an alloy within the composition range
15/17 per cent molybdenum - 6/8 per cent chromium - 4/6 per cent iron -
0.04/0.,08 per cent carbon - balance nickel, designated as INOR-8, to be
an attractive structural material for use in a nuclear power reactor
fueled with molten-uranium-bearing fluoride salts.

The present study enlarges upon the technology of the alloy
INOR-8 through an investigation of the influence of camposition vari-
ation upon the 1500°F creep~rupture strength and microstructure of
alloys encompassed by the range 10/20 per cent molybdenum - 5/10 per
cent chramium - 4/10 per cent iron - 0.5 per cent aluminum - 0.5 per
cent manganese - 0.06 per cent carbon - balance nickel. The campo-
sition of the individual alloys was varied systematically with the
intent that by direct comparison the effect of an element upon strength
could be determined. All alloys were tested in creep-rupture at a
stress of 10,000 psi in the annealed condition. The criteria used to
evaluate the strength of the alloys were the times requiréd to reach
strains between 1 and 10 per cent.

The results could not be explained in gimple terms of compo-
sition variation since the principal factors affecting the strength

of the alloys were: solid-solution elements, carbide and non-carbide

 
O

aging reactions, the presence of M6C-type carbides in the micro-
structures, and grain size.

From the standpoint of their creep-rupture strength, it was
possible to conveniently group the alloys according to the three con-
centrations of molybdenum studied: 10, 15, and 20 per cent. It could
be concluded from the analyses and microstructures of these alloys that
the relative strength contribution of each of the previously mentioned
factors varied between the individual groups.

The cambined effects of solid-solution strengthening by molyb-
denum and the increase in quantity of dispersed M6C—type carbides which
this element pramoted in the annealed materials were the predominant
factors which increased the strength of the alloys grouped by molyb-
denum content. The only exception noted was in the case of the 20 per
cent molybdenum ~ 7 per cent chromium - 10 per cent iron alloy which
precipitated a non-carbide phase as a consequence of crossing a new
phase boundary. This phase contributed noticeably to creep-rupture
strength in the later stages of test.

The contribution of chromium and iron to the strength of the
alloys within the individual groups could not be established with
certainty due to simultaneous variations in other factors affecting
creep-rupture behavior.

To obtain a better indication of the strengthening influence of
chromium and iron, creep-rupture studies were conducted on low-carbon
"high-purity™ alloys. Although an analysis of the data was complicated

by the presence of a limited amount of carbide precipitation and by
grain-size variations, the influence of chromium was found to be
significant when 5 to 10 per cent was added to the 15 per cent
molybdenum - balance nickel base. However, the presence of 10 per
cent chromium in the base composition showed the most pronounced
strengthening influence. The strengthening effect of iron was inter-
preted as being insignificant when amounts up to 10 per cent were
added to the 15 per cent molybdenum - 7 per cent chromium - balance
nickel base.

A general consideration of all data obtained fraom this investi-
gation favorably supports the composition specification placed upon the

alloy INOR-8.

 
CHAPTER 11
INTROTUCTION

The advancement in the technology of a nuclear power reactor
fueled with molten~uranium-bearing fluoride salts has been concerned,
in part, with the development of a structural material which will be
compatible with the reactor operating conditions. In such a reactor
the material would be subjected to several corrogive environments at
elevated temperatures in addition to complex stresses derived from
flowing fluids, temperature gradients, and thermal cycles; conse-
gquently, it is necessary that the material meet rigid requirements,
among them being:

1. sufficlent strength and reasonable ductility at elevated
temperatures;

2. good corrosion reslistance to the molten fluoride salts;

3. good oxidation reszistances;

4, favorable fabricability for the production of a variety
of shapes; 1.e., plate, sheet, bar, wire, tubing, etc.;

5. good weldabllity and brazeability; and

6. suitable nuclear properties.

It has been found through various test programs conducted over
the past several years, that of the commercially available alloys,

‘X'
Inconel (80 Ni - 14 Cr - 6 Fe) ard Hastelloy B (67 Ni - 28 Mo ~ 5 Fe)

 

*
All alloy campositions are expressed in weight per cent.

 
 

vere the most pramising for this application; however, neither is an
"ideal" material with respect to the above requirements. With parti-
cular regard to the elevated temperature strength of the two alloys,
Hastelloy B is much superior to Inconel at 1500°F, as shown in
Figure 1. In structures fabricated from Inconel, specification of
design stresses must take into account the deterioration of its
strength due to a significant amount of corrosion by the molten
fluoride salts as well as that due to temperature. Thus, in many
cases, Inconel becames a marginal alloy where thin sections are en-
countered. Hastelloy B, on the other hand, is virtually unaffected
by the molten fluoride salts and, at the same time, possesses high
strength; however, extended service in the temperature range of approxi-
mately 1200 — 1600°F results in a decrease in its ductility to such a
degree that it also becames a marginal material.

A recent study of the aging characteristics of Hastelloy B by
R. E. Clausing, et al,l has shown the most prominent loss in ductility
of the alloy to occur after aging at 1300°F, due to the precipitation
of an intermetallic compound interpreted as being the beta phase
(NiuMo) of the nickel-molybdenum system. Between 1500°F and 1650°F,
the precipitate observed upon aging was different fram that noted at
1300°F and was tentatively identified as the gamma phése (NisMo).
'The effect of the gamma phase on decreasing the tensile ductility of
Hastelloy B was not quite so marked as that of the lower temperature
beta phase. The assumed intermetallic campounds were deduced fram the

nickel-molybdenum equilibrium diagram2 shown in Figure 2, 1t belng

 
 

UNCL ASSIFIED
ORNL-L.R-DWG, 26319

 

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STRESS (psi)

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TIME FOR FAILURE (hr}

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Figure 1. Comparison of the stress-rupture properties of Inconel and Hastelloy B at 1500°F.

 
 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

UNCLASSIFIED
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Weight Percentage Molybdenum
Figure 2. Nickel-Molybdenum equilibrium diagram.

 
recognized that the presence of iron and other elements in Hastelloy B
could have an influence on the location of the phase boundaries.

In view of the over-all requirements of the "ideal" alloy, it
became apparent that these could best be met by a material having the
desirable features of Inconel and Hastelloy B. During the past few
years, an alloy development program at the Oak Ridge National
laboratory has been concerned with the evaluation of numerous nickel-
base alloys with a primary strengthening addition of 15 — 20 per cent
molybdenum. This amount of molybdenum is within the solubility limits
of molybdenum in nickel at elevated temperatures, and therefore, em-
brittlement associated with the precipifafibfi of l\TiXMoy intermetallic
compounds does not occur. BSubsequent additions made to the nickel-
molybdenum base coamposition to augment its properties included such
elements as chromium, iron, niobium, vanadium, tungsten, aluminum,
titanium, end carbon. Data obtéined fram screening tests designed to
investigate strength, fluoride-salt corrosion, oxidation, fabricability,
and weldability of the various alloys showed the necessity of a com=
positiofi balance since enhancement of a particular property was obtained
at the expense of one or more other properties. At the present time,
the best compromise material, in light of the origifial requirements of
the "ideal" alloy, is the camposition, 15/17 Mo -~ 6/8 Cr = 4/6 Fe -
0.04/0.08 C ~ balance Ni, designated as INOR-8. This alloy is classi-
fied as the solid=-solution type with the exception, however, of carbide

particles which are stable over a wide temperature range.

 
 

During the course of the alloy screening tests the potential
importance of a number of variables became“evident. Among those
variables incompletely understood were the influence of heat treatment,
melting practice, and changes in composition upon the strength of a
given alloy. Of these three subjects, it was believed that a program
set up to investigate the influence of composition variation upon the
strength, as well as the microstructure, of alloys within a range
encompassing that of the alloy INOR-O would prove the most beneficial.
Such an investigation was carried out and is described in the following

chapters.

 
 

-10-

CHAPTER III
OBJECTIVE

The objective of thls investigation was to determine the
influence of composition variation upon the 1500°F creep-rupture
strength and microétructure of molybdefium-chromium-iron-nickel base
alloys. The alloys fof this study were encompassed by the range
10/20 Mo - 5/10 Cr - M/lO Fe - 0,5 AL - 0.5 Mn - 0.06 C - balance Ni.
The composition of the individual alloyé was varied systematically with
the intent that by direct comparison, the effect of an element upon

strength could be determined.

 
CHAPTER 1V

EXPERIMENTAT, PROCEDURE

Two general serieg (I and II) of alloys were prepared for this
investigation. Series I was prepared to show the influence of vari-
ations in the molybdenum, chromium, and iron contents at a constant
carbon content on the creep-rupture behavior and microstructure of the
resultant alloys. The melting stock used for preparing the alloys of
this series was of the following reported purity: nickel pellets,
99.9 per cent; sintered molybdenum bar ends, 99.8 per cent; alumino-
thermic chromium, 99.3 per cent with approximately 1000 ppm oxygen;
vacuun-melted ingot iron, 99.9 per cent; and carbon in the form of
graphite. Aluminum and manganese were added as malleableizing agents.
Previous work with nickel-molybdenum=chromium alloys at ORNL has shown
them to be subject to cracking during hot-rolling in the absence of
malleableizers, even though the alloys were prepared by vacuum-melting
techniques from reportedly good-quality melting stock. It is believed
that the hot-cracking tendencies are caused, principally, by the
residual gases present in the materials.

The presence of carbon in the alloys of series I produced creep-
rupture results which indicated that carbon through carbide formation
was masking the strengthening influence of chromium and iron. This
fact prompted the preparation of the alloys of series II, The alloys
of series II are referred to as "high-purity" alloys and were prepared

from the same nickel and iron melting stock as the slloys of series I

 

 
-12-

however, arc-cast molybdenum and special high-purity chromium flakes
(lhO ppm oxygen) were substituted for the sintered molybdenum bar ends
and alumino-thermic chromium, respectively. No intentional carbon or

- malleableizing agents were added to the alloys of series II.

Melting and Casting

All alloys were prepared by vacuum-induction-melting. A total
charge of 1450 grams per alloy was placed in a zirconia crucible, out-
gassed by intermittent application of power, and finally melted under a
vacuum of approximately 100 microns of mercury. Each alloy was held in
the molten state from twenty to thirty minutes tc insure solution of the
charge, and then cast into a graphite mold to form an ingot one and one-

half inches in diameter by four inches in length, excluding the hot-top.

Ingot Analysis

The hot-tops were cut fram the ingots and turnings for chemical
analysis were taken from the bottam of each hot-top across the entire
transverse section. A prior skin cut, which was discarded, was made on
the outside diameter of the hot-tops before collecting the turnings.

The results of the analyses performed by the Analytical Chemistry
Division of ORNL are shown in Table I. The naminal composition of each
alloy based upon the weight of the elements making up the charge is also
given in Table I for comparison. In general, the nominal compositions

were in good agreement with the analyzed compositions.

 
TABLE T

NOMINATL AND ANALYZED COMPOSITIONS OF THE ALLOYS

 

 

 

 

 

 

 

 

 

 

Allo Nominal Composition (Wt %) Anslyzed Camposition (Wt %)
- Ne. Ni Mo Cr Fe Al Mn Cc Ni Mo Cr Fe Al Mn C
Beries I VI-43 Bal 10 S5 L4 0.5 0.5 0.06 Bal 9.87 L.94  L4.00 0.79" 0.59 0.076
VI-44 Bal 10 5 10 0.5 0.5 0.06 Bal 9.83 L4.86 9.63 0.90 0.61 Q.06
VI-47 Bzl 10 7 4 0.5 0.5 0.06 Bal 9.61 6.59 3.76 0.81 0.59 0.07
VI-48 Bal 10 7 10 0.5 0.5 0.06 Bal 9.42 6.58 9.29 0.82 0.53 0.07
VI-53 Bal 10 10 4 0.5 0.5 0.06 Bal 9.62 10.01 4.19 0.78 0.50 0.064
VI-54 Bal 10 10 10 0.5 0.5 0.06 Bal 10.93 9.71 10.67 0.88 0.53 0.068
VI-45 Bal 15 5 L 0.5 0.5 0.06 Bal 14,10 L.79 3.93 0.77 0.55 0.073
VI-46 Bal 15 5 10 0.5 0.5 0.06 Bal 16.27 4,90 10.33 0.81 0.57 0.090
VI-49 Bal 15 7 L4 0.5 0.5 0.06 Bal 15.50 6.83 L4.32 0.88 0.59 0.075
VI-50 Bal 15 7 10 0.5 0.5 0.06 Bal 14.37 6.99 10.21 0.86 0.64 0.081
VP-55 Bal 15 10 L4 0.5 0.5 0.06 Bal 15.94% 9.80 L4.25 0.81. 0.53 0.077
VI-56 Bal 15 10 10 0.5 0.5 0.06 Bal 15.76 9.84 10.29 0.89 0.43 0.077
VI-57 Bal 20 5 L 0.5 0.5 0.06 Bal 21.06 L4.88 L.31 0.86 0.56 0.07
VI-58 Bal 20 5 10 0.5 0.5 0.06 Bal 19.72 7.19 10.38 0.81 0.41 0.069
VI-59 Bal 20 7 L4 0.5 0.5 0.06 Bal 18.60 6.93 L4.23 0.94% 0.41 0.075
VI-60 Bal 20 7 10 0.5 Q.5 0.06 Bal 20.70 7.11 10.17 0.76 0.38 0.066
Series II VI-Q0 Bal 15 - - - - - Bal 14.34 - - - - 0.022
- Vr-89 Bal 15 3 - - - - Bal 14.59 2.98 - - - 0.018
Vr-88 Bal 15 5 - - - - Bal 14.39 5.04 - - - 0.024
Vvr-87 Bl 15 7 - - - - Bal 14.89 7.03 - - - 0.025
VI-86 Bal 15 10 - - - - Bal 14.89 10.19 - - - 0.017
VI-9l Bal 15 7 L - - - Bal 15.53 5.0 2.05 - - 0.010
VI-92 Bl 15 7 7 - - - Bl 15.24 7.19  7.19 - - 0.021
VI-93 Bal 15 7 10 - - - Bal 14.79 7.16 10.18 - - 0,024

 

-ET"'

 
Ingot Fabrication

The rough, as-cast surface of each ingot was ground smooth
prior to break-down accamplished by hot-rolling in air at a furnace
temperature of 2l50°F3 Reductions of 50 mils in thickness per pass
were given from_the initial one and one-hglf inch diameter to 0.750
inch thick. From a thickness of 0.750 inch to 0.275 inch reductions
of 30 mils in thickness per pass were giveno After hot-rolling, the
alloy strip; were pickled in a hot aqueous solu;.tion3 of 10 per cent
sulphuric acid containing 5 per cent by weight of sodium nitrate and
5 per cent by weight of sodium chlorideo‘ At this stage the materials
were examined visually and edge and surface cracks which developed
during hotwrollingwére ground out. The‘strips were subsequently cold-
rolled to 0.065 inch thick at a reduction schedule of approximately
3 mils in thickness per pass. Although alloys of this type work-harden‘
quite rapidly, it was possible to cold-roll fram 30 to 4O per cent in
thickness before an intermedigte annealing treatment for one-half‘hour
at 2150°F in a hydrogen atmosphereo The above procedure was followed
as nearly as possible for all the alloys.

In genera;, the fabricabllity of both series of alloys as
determined byrthé extent of cracking during hot-rolling was satis-
factory with one exception being an alloy of series I, VT~60 (20 Mo -
7 Cr - 10 Fe - 0.5 AL - 0.5 Mn - 0.06 C - balance Ni). Strip suitgble
for creep-rupture specimens could not be obtained due to excessive

cracking.
 

-15-

Creep~Rupture Testing

The creep-rupture tests were carried out in the Mechanical
Properties Laboratory of the ORNL Metallurgy Division. A description
of the creep~testing facilities of this ILaboratory has been reported
previously by D. A. Douglas and W. D. Manly.u Details from this report
which were pertinent to this investigation, i.e., description of appa-
ratus, are presented 1n the Appendix.

After stress-relieving the 0.065-inch-thick strip for one-half
hour at 1600°F, two sheet=-type creep~rupture specimens, as illustrated
in Figure 3, were machined fram each alloy. Prior to test, all
specimens were annealed at 2100°F for one-~half hour in a hydrogen
atmosphere followed by cooling in the furnace cold-zone considered
equivalent to an air-cool. Each specimen was tested in creep-rupture
at a stress of 10,000 psi (constant load), a temperature of 1500°F, and
in an inert atmosphere of argon. After fixing a specimen in a creep
frame, the temperature was brought up to 1500°F. The control tempers-
ture during test as well as the temperature gradient along the specimen
gage length was maintained within % 5°F. Loading of the specimen tock
place immediately upon achieving a uniform temperature along the gage
length.

For the alloys of series I, microscopic and dial-gage extension
readings were taken every two hours during the first eight hours of the
test, and thereafter, readings were made once every twenty-four hours
until rupture occurred. Extension readings were made on the "high-

purity” alloys of series II every one-half hour during the initial

 
 

 

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UNCL ASSIFIED
ORNL-LR-DWG. 14674

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Figure 3. Sheet-type creep-rupture specimen.

 

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~17-

stages of test to closely follow the strain vs time curves to 1 per
cent strain. Thereafter, readingé were taken on these alloys every
two hours. After a specimen ruptured, it was allowed to furnace-cool
and a total elongation measurement was made.

The criteria used to evaluate the strength of the alloys were
times to 1, 2, 5, and ld.per cent strain. Because more freguent
extension readings were made on the "high-purity” alloys during the
initial stages of te;t, the time to 0.5 per cent strain was also used
as a strength criterion for this series of élloys.

In order to make a relative comparison of the strengths of the
alloys, a three and one~half inch gage length was arbltrarily selected
for camputing the per cent strain from the extension measurements. The

point of zero elongation was taken as the reading obtained immediately

after the application of the full load on the specimen.

Grain Size Measurements

The standard annealing treatment given each creep-rupture
specimen prior to test produced different grain sizes between the
various alloys, thus introducing another variable to be considered in
evaluating the test results. To obtain an indication of the wvariation,
grain-size measurements were made on a longitudinal section of an
unstressed end of one'creep-rupture specimen per alloy using the Heyn
‘Procedure.5 Briefly, the method consists of counting the number of

gralns in a magnified image of the specimen intersecting a line of

 
~18-

known length. By dividing the length of the line by the product of the
number of intersecting grains times the magnification, a value of the
average graln diameter is obtalned. Each reported grain size repre-
sents an average of ten readings taken at different locations in a

given sample.

Aging Studies

All alloys of series I were investigated for their aging res-
ponse. Coupons were cut fram the 0.065-inch-thick strip, annealed one-
half hour at 2100°F in a hydrogen atmosphere and cooled in the furnace
cold zone. AllL coupons were placed in a quartz tube,sealed off under
a vacuum of 0.l micron of mercury, and then aged at 1500°F for 5, 25,
50, 100, and 1000 hours. All aging heat-treatments were carried out
in Kanthal-wound furnaces which were constructed at ORNL. Chromel-
alumel thermocouples were used for controlling and recording tempera-
ture. Periodically, furnace temperatures were checked with a
standardized platinum/platinum-10 per cent rhodium thermocouple and
a Rubicorn Potentiometer. Figure 4 shows a photograph of the heat-
treating furnaces with their temperature-controlling and recording
instruments. Upon coampletion of all aging treatments, the capsules
were crushed under water to effect rapid cooling of the coupons.

A surface of each coupon parallel to the rolling direction was
prepared for metallographic examination using the procedure to be

described later. The aging response of the alloys was determined by

 
 

 

 

Figure L.

UNCL ASSIF [ED
PR P HO TO 41669

 

Heat-treating furnaces with temperature-controlling and recording instruments.

|-

O

 
-20-

hardness measurements made on these specimens as well as a solution-
annealed standard of each alloy. The average hardness reported for s
given specimen represents an average of four measurements taken at |
different locations on the specimen. All hardness measurements were
made with a Wilson Tukon Micro-Hardness Tester using a 10 kilogram
load with a 16 millimeter objective and a 136° diamond pyramid
indenter. | | |

Because carbide precipitation was not anticipated in the "high-
purity" alloys of series II, but did occur to a limited extent as will
be pointed out in Chapter V, one coupon of each alloy of this series-
was aged for lOO hdurs at 1500°F to determine its "equilibrium"
structure; fiardfiess and metallographié studies were performed on

these specimens and a solution-anneasled standard.

Decarburization Studies

In order to present evidence that certain observed precipitétes
were carbide particles, selected ccmpositions were subjected to a
decarbfirization treatment and examined for the disappearance of precipi-
tates. The treatment conéisted of rolling alloy strip to 0.012 inch.
thick, heat-treating for 100 hours at 2200°F in a hydrogen atmosphere
to effect decarburization, dropping the furnace temperature to 1500°F
and aging for an additional lOO hours, followed by water-quenching from

temperature. The microstructures of the alloys were examined and com-

pared with the microstructures of the same alloys after aging at 1500°F

 

 
-21-

in an evacuated quartz capsule. A carbon analysis was also obtained on
the decarburized materials.

The alloys subjected to this treatment included: 10 and 15 per
cent molybdenum alloys of series I with the lowest and highest chramium
plus iron contents; 20 per cent molybdenum alloys of series I con-
taining 5 per cent chramium plus 4 and 10 per cent ironj and the “high-
purity" alloys of series II containing 7 per cent chramium plus O, 7,

and 10 per cent iron.

Metallographic Studies and Procedures

In addition to a solution-annealed coupon of each alloy, the
alloy specimens aged for carbide precipitation at 1500°F, and those
subjected to the decarburization treatment, metallographic studies
were conducted on an as-cast specimen taken from the ingot hot-top of
each alloy, and longitudinal and transverse sections in both the gage
length and grip of one creep-rupture specimen per alloy.

All speclmens which were prepared for'metallografihic examination
were first mounted in the conventional manner in bakelite.‘ Initial
grinding was done on lead laPS‘employing'American Optical 302, 303-1/2,
and 305 corrundum abrasive (listed in order of decreasing particle
size), followed by intermediate and final polishing on a Syntron Vibro-
Polisher. Three steps were used for intermediate and final polisHing:
(1) a silk cloth with 0.3 micron aluminum oxide abrasive (ILinde A),

(2) a micro cloth with linde A abrasive, and (3) a micro cloth with

 
w2

0.1 micron aluminum -oxide abrasive (Linde B). After polishing, all
specimens were etched in glycera regia consisting of 1 part HNOS,
3 parts HCl, and & parts glycerine. Photamicrographs were made with

a Bausch and Lamb Research Metallograph using bright fleld illumination.
 

-23-

CHAPTER V
RESULTS AND DISCUSSION
CREEP-RUPTURE STUDIES

The interpretation of the creep-rupture data obtained from the
alloys studied for this investigetion could not adequately be made in
gimple terms of camposition variations. This complication arose since
camposition variations caused not only solid-solution strengthening,
but eleo, variations in grain size, dispersed particles, and precipi-
tation reactions within the materials. Analysis of the data showed
that these factors, which are known to affect the creep-rupture be-
havior of an alloy, were interrelated to varylng degrees in establishing
the properties of the alloys. The observed creep-rupture behavior,
therefore, is the resultant of these combined variables.

The carbon intentionally added to the alloys of series I was
very effective in introducing into these materials changes in the micro-
structure which affect creep-rupture behavior. For illustration,
Figures 5, 6, and 7 show the as-cast microstructures of three alloys of
series I at the different molybdenum contents with 7 per cent chromium -
4 per cent iron - 0.5 per cent aluminum - 0.5 per cent manganese -

0.06 per cent carbon - balance nickel. In addition to the face~centered-
cuble matrix, sll contained at least one additionsl phase which formed
by a eutectlc decamposition in the 15 and 20 per cent molybdenum alloys

and which apparently precipitated from solid solution In the 10 per

cent molybdenum alloy. In contrast, the as-cast microstructure of a

 

 
UNCLASSIFIED
Y-2484 1

 

Figure 5. Alloy VI-47, 10 Mo - 7T Cr - 4 Fe - 0.5 AL - 0.5 Mo -
0.06 ¢ - Balance Ni. As-cast. BEtchant: Glycera Regia. 500X.

\_'_:‘u £y g -" i a ','!'..-l,‘ "(r;
. My e :r,«r XF UNCLASSIFIED
“\f';;_[/ Bl fl\" Y-24842

 

Figure 6. Alloy VI-49, 15 Mo - 7 Cr - 4 Fe - 0.5 AL - 0.5 Mn -
0.06 C - Balance Ni. As-cast. Etchant: Glycera Regia. 3S00X.
 

UNCL ASSIF IED

Y-24843

 

Figure 7. Alloy VI-59, 20 Mo - 7T Cr - 4 Fe - 0.5 AL - 0.5 Mn -
0.06 C - Balance Ni. As-cast. Etchant: Glycera Regia. 500X.

\ UNCLASSIFIED
L& Y-24844

Figure 8. Alloy VI-91, 15 Mo - 7 Cr - 4 Fe - Balance Ni.
As-cast. Etchant: Glycera Regia. 500X.
-26 -

low-carbon "high~purity” alloy of series II with 15 per cent
molybdenum - 7 per cent chromium - 4 per cent iron - balance nickel

is shown 1n Figure 8. The cleanliness of this microstructure indicates
that the second phase observed in the 0.06 per cent carbon alloys was a
carbide.

Microstructures developed in these alloys after rolling the
ingots to 0.065-inch-thick strip, and solution-annealing specimens of
each at 2100°F for one-half hour in a hydrogen atmosphere are shown in
Figures 9, 10, 11, and 12, By camparing Figures 9, 10, and 1l, a trend
was established with respect to an incremsing quantity of globular M6C~
type* carbides present as stringers in the materials with an igcreasing
molybdenum-to-carbon ratio. The presence of these carbides influenced
the as-solution annealed grain size of the series I alloys with the
result that it ranged from coarse to fine, depending upon the total
quantity of particles present. Furthermore, a fine-grain-boundary
precipitate was found in the alloys with 10 and 15 per cent molybdenum
vhich was presumably a carbide formed during cooling of the specimens
to room temperature., The fine-grain-boundary precipitate was for the
most part absent in the 20 per cent molybdenum alloys which reflected a

tie-up of most of the total carbon in the globular M6C—type carbldes.

 

*fiecent Work6 performed by the Westinghouse Electric Corporation
for ORNL has shown the microstructure of an as-forged sample of the
alloy INOR-8 with 0.13 per cent carbon to contain large globular parti-
cles. These particles were ldentified as an MgC-type carbide by x-ray
diffraction analysis of the residue obtained from electrolytic dis-
solution of the sample. X-ray spectrographic analysis of this residue
showed molybdenum to be the principal constituent together with nickel.
The MGC carbide was, therefore, considered to be (M0§N1)6Co

 
   

__UNCLASSIF IED
: 822

Figure 9. Alloy VI-47, 10 Mo - T Cr - 4 Fe - 0.5 AL - 0.5 Mn
0.06 C - Balance Ni. Solution-annealed 1/2 hr. at 2100°F. Etchant:
Glycera Regia. 100X.

T R IO T e g B BRI ek T

Remige TN 0T el TR e Y " TTUNCL ASSIFIED

gl v 3 K, - ¥ "h.l‘: 5 g - _-"'-- . :
NSt ey

W 3 ReE \ 3 T L "
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2k TR e Q’J—-’-*T: i . JJT[FJ o '-ET:-' -y gt S =s A ST
Th of T Fggef sty |

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= ’ @ Sl e L SIS o= ey

e -fi:ilc'-l"-r' S ey e e, Ty e b E:lq.- ‘ F e a L
D e et 1 o e - N N gy P e = A ey

Figure 10, Alloy VT-49, 15 Mo - 7 Cr - &k Fe - 0.5 Al - 0.5 Mn -
0.06 C - Balance Ni. Solution-amnealed 1/2 hr. at 2100°F. Etchant
Glycera Regia. 100X.

 
Y-24926

    

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- iy o i F ST =gl e S T
e iy Py -
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e - -? & g T e
- 5 Tt - — -
b _-,-I - -..._,_,_::E;__:h“. b __"_—-,. . hfl.m__..‘” b
- e .

Figure 11. Alloy VI-59, 20 Mo - T Cr - 4 Fe - 0.5 AL - 0.5 Mn
0.06 C - Balance Ni. Solution-annealed 1/2 hr. at 2100°F. Etchant:
Glycera Regia. 100X.

\ T 1/ '_UNCLASSIFI ED
%jfi” \\.J< i R

e — |
|

L e 2
N e | S =
N =5
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Figure 12. Alloy VI-91, 15 Mo - 7 Cr - 4 Fe - Balance Ni.
Solution-annealed 1/2 hr. at 2100°F. Etchant: Glycera Regla. 100X.
 

The hydrogen atmosphere used during solution=annealing effective-
ly caused decarburization around the edges of the 0.06 per cent carbon
alloys, particularly at the 10 and 15 per cent molybdenum concen-
trations resulting in a graine-size gradient in these alloys. Grain-
slze measurements reported for the alloys, however, included the
average of all grains across the thickness of a glven specimen.

A presentation of the creep-rupture data which were obtained on
the various alloys follows 1n this chapter, and are discussed in terms

of the several variables which were encountered.

10.- Per Cent Molybdenum Alloys with 0.5 Per Cent Aluminum - O.5 Per Cent

 

Menganese - 0.06 Per Cent Carbon and Varying Percentages of Chramium

 

and Iron

A summary of the creep-rupture data obtained gn the 10 per cent
molybdenum alloys of series I is presented in Table II and illustrated
by the graph in Figure 13. Bach creep~rupture test has been repre-
sented by a bar in Figure 13 on which is indicated at the appropriate
times, stralns of 1, 2, 5, and 10 per cent as well as the specimen
elongation at fracture. The times to the specified values of strain
vere interpolated fram a plot of strain vs time for each test.

Obvious metallurgical factors which could have Influenced the
measured creep rate of this group of alloys were: a varimation in grain
size between the individual alloys, carblde precipitation occurring
within the alloys at 1500°F, and variations in the chromium and iron
contents. 1In addition, a certain amount of scatter existed between

the data for the duplicate tests of any one allioy.

 
 

TABLE IT

SUMMARY OF CREEP-RUPTURE DATA ON Ni-BASE ALLOYS CONTAINING 10 Mo - 0.5 Al -
0.5 Mn - 0.06 C AND VARYING PERCENTAGES OF Cr AND Fe
Test Conditions

Temp ; 1500°F
Stress: 10,000 psi
Atmos : Argon

 

 

 

Time to Specified

 

 

Camposition (a) Ruptiure Elong- Average
Test Alloy Variable (Wt %) Strain >/ (gr.) Life ation Grain Dia.
No. No. Gr Fe 1% 2% 5% 10% (Hr.) (%) (rum )
10-8-1. flfiTmu3 5 L 1k.5 27 N 155' 233.7 31.25 .
10-8-11 1k 28.5 63 105 156.9 28.57 0.0833
10-8-2 VP-44 5 10 14,5 27 62 100 151.6 20. 54 -
10-8-20 16 28 54 72 75.8 1.3.39 0.0653
10-8-5 VP47 7 L 13 23 53 83(b) 8h.7 10.71 O0.1h47
10-8-8 16 28 59 - 8L,7 8.93 -
10-8-6 VE-48 T 10 1.5 26 55 90 106.2 16.07 -
10-8-19 19 33 63 98 131.8 16.07 0.0598
10-8-21 VEP-53 10 Y 23.5 43 81 . (p) 99,7 7.1k 0.0833
10-8-29 19 33.5 64 - 71.5 6.25 -
10-8-15 VT-54 10 10 16.5 31 68 118 210.9 37.5 0.143
10-8-31 | 20 36 73 117 134.3 13.39 -

 

(a) Microscope readings except where noted.
(b) Dial-gage readings.

.
om )
UNCL ASSIFIED
ORNL-LR-DWG. 26324

NOMINAL COMPOSITION (wt %) ———————  RUPTURE: 37;50 hr
BASE: 10 Mo—05 Al— 05 Mn—0.06 C—BALANCE Ni — ELONG. 12%

; 5Cr 5 Cr T Cr 7 Cr 10 Cr 10 Cr
VARIABLE: | 4 e 10 Fe 4 Fe 10 Fe 4 Fe 10 Fe

400 TEST CONDITIONS
TEMP: 1500°F

STRESS: 10,000 psi
ATMOS: ARGON

200 *DIAL GAGE STRAIN MEASUREMENTS

1000
800

600

100
80

60

40

20

TIME TO SPECIFIED STRAIN (hr]

 

1
ALLOY: INCONEL VT-43 vT—-44 vT-47 vT-48 VT-53 VT-54 HASTELLOY B

Figure 13. Bar-graph of creep-rupture test results obtained at 1500°F, 10,000 psi.,
on Ni-base alloys with 10 Mo - 0.5 Al - 0.5 Mn - 0.06 C and varying percentages of Cr and Fe.

 

T¢

 
The results of average grain-diameter measurements on an un-
stresséd portion of the creep-rupture specimens are shown in Table II.
A reason for the variation in gralin size was found in that greater
numbers of stringers produced finer grain sizes. The amount of carbide
stringering, however, could not be correlated well with the carbon
analyses of the alloys, This lack of correlation was attributed to
the possibllity of carbide segregation being present in the alloys, or
that the annealing'treatments given the alloys during ingot fabrication
caused decarburization. If either or both of these explanations be
valid, then the carbon content of an individual sample was not repre-
sentative of the analyzed carbon content of the parent ingot.

The alloys of this group were all found to be subject to carbide
precipitation. Table IIT shows the range and average hardness of the
various coupons of each alloy after solution-annegling at ElOO?F.and
subsequent aging at 1500°F, Figure 14 1s a graphical representation of
the average hardness of each alloy as a function of time at 1500°F,
campared with its hardness in the solution-annealed condition. In
general, pesk hardness of the alloys was reached after 25 to 50 hours
at temperature. It should be mentioned that some scatter was encoun-
tered in the hardness values obtalned on the various coupons. This
could be attributed to orientation effects accentuated by large grain
size as well as segregation effects of carbide bamding in the rolling
direc£ion°

A general chrelation could be made between the hardness change

upon aging at 1500°F and the resultant microstructures. To illustrate

 
TABLE ITI

DIAMOND PYRAMID HARDNESS DATA ON Ni-BASE ALLOYS CONTAINING 10 Mo -
0.5 Al - 0.5 Mn - 0.06 C AND VARYING PERCENTAGES OF Cr AND Fe.
SOLUTION -ANNEALED 1/2 HR. AT 2100°F, AGED AT 1500°F.

 

 

 

 

 

 

ot ol Age st 25007

Alloy (Wt. %) at 2L00°F 5 Hr. 25 Hr. 50 Hr. 100 Hr. 1000 Hr.

No. Cr Fe  Range Avg. Range Avg. Range Avg. Range Avg. Range Avg. Range Avg.
VT-43 5 L 138141 140 178189 184 188-196 193 175-191 183 169-178 17h 1h3-a52 147
VT-4L4 5 10 135151 145 158473 166 162-175 168 163-172 168 149161 157 147-155 150
VT-47 7 L 130146 139 157166 162 157-168 163 161-165 163 154163 159 1Lh3-251 146
VT-48 T 10 138-145 141 161164 162 157-159 158 149-158 154 150161 157 136-146 143
VT-53 10 4y 132-1hk9 1L1 172177 175 173~l76 174 170-180 176 169176 172 156—-170 16k
Vr-54 10 10 124440 135 160-165 163 161-168 163 166-186 174 156157 157 143151 148

 

..EE‘..

 
 

UNCL ASSIFIED
ORNL+L R-DWG, 26320

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

210
NOMINAL COMPOSITION (wt %)
BASE: 10 Mo—0.5 Al—0.5Mn—0.06C —BALANCE Ni
200
] S Cr 5Cr 7 Cr 7Cr 10 Cr 10 Cr
VARIABLE:| 5 £ 10 Fe 4 Fe 10 Fe 4 Fe 10Fe ||
[
{90
v [ -—
i
Z 180
O
[0
<l A
I
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§ |
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a 1
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150
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_E o E E -E - E = E - E —
) < 'e) — njL f. o < 3 ] o< — n|< -
| = l<|o VI —1 |@I<| . OlE| RN
140 NSl El . ol<| = ol<|x of<|E ol €| = NolEl s
wlo|< 2 nlo|= Blo|t ©Olo|s : 0o 0lo
AR B Rgl LR K Re 2| |78 8I
2 Q 2 oz o o . Q
<1 <l =
Z
430 :
ALLOY: VT-43 VT-44 vT-47 vT-48 VT-53 VT-54

Figure 14. Bar-graph of average hardness of Ni-base alloys with 10 Mo - 0.5 AL -
0.5 Mn - 0.06 C and varying percentages of Cr and Fe. Solution-annealed 1/2 hr. at 2100°F,

aged for 5, 25, 50, 100, and 1000 hr. at 1500°F.

1t
-35-

this, Figure 15 shows a series of photomicrographs taken of alloy VT-43
(10 Mo - 5Cr - 4 Fe - 0,5 AL - 0.5 Mn - 0,06 C - balance Ni) during
the course of aging in the absence of mechanical stress. In the
solution-annealed condition, the microstructure contained a fine-grain-
boundary precipitate and a few stringers of larger M6C-type carbides.
After five hours at 1500°F, a matrix precipitate was in the early stages
of development with fine particles being precipitated along preferred
crystallographic planes. The progress of precipitation continued with
time when, after 50 hours, discrete particles were observed which
showed only slight change in size up to 1000 hours at temperature. The
microstructure of the gage length of a creep-rupture specimen of alloy
VI-43 is shown in Figure 16 and was found to be identical to that
observed in the unstressed coupons aged for 50 hours or more.

It was slignificant to note that the aging response of alloy
VI=43 was greater in terms of pesk hardness and guantity of precipi-
tated carbides than that shown by the other alloys of this group. An
apparent correlation existed between this fact and the reported carbon
content of 0.076 per cent for alloy VI-43 which was higher than that
reported for the other campositions.

The least amount of aging, noted metallographic&lly‘in the 10 per
cent molybdenum alloys was found in alloy VI-54 (10 Mo - 10 Cr - 10 Fe -
0.5 AL -~ 0.5 Mn - 0.06 C - balance Ni). Figure 17 shows typical micro-
structures exfiibited by this alloy upon aging. During the early stages
there was evidence of previpitation of fine carbides within the grains,

but longer times at temperature promoted a concentration of these

 
 

UNCLASSIFIED ! UNCL ASS[;:IED

   

UNCLASSIFIED
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Solution Annealed. DPH 140 5 Hr. at 1500°F. DPH 184 25 Hr. at 1500°F. DPH 193
UNCL ASSIFIED UNCL ASSIFIED i - —  UNCLASSIFIED
Y. 24830 Y-24B832 Y¥-23967 L
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Figure 15. Alloy VI-43, 10 Mo - 5Cr - 4 Pe - 0.5 Al - 0.5 Mn - 0.06 C - Balance Ni.
Solution-annealed 1/2 hr. at 2100°F, aged at 1500°F. Etchant: Glycera Regia. 500X.

 
37

 

Figure 16. Alloy VT-43, 10 Mo - 5 Cr - 4 Fe - 0.5 Al -
0.5 Mn - 0.06 C - Balance Ni. Iongitudinal section in gage length of
creep-rupture specimen. Rupture life: 156.9 hr. Etchant: Glycera
Regia. 500X.

 
/ UNCL ASSIF IED i UNCLASSIFIED UNCL ASSIFIED
/ Y-23485 ! Y-249 18 __( Y-249 19

 

\ [ ~ /
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2t / ‘
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Solution Annesled. DPH 135 5 Hr. at 1500°F. DPH 163 25 Hr. at 1500°F. DPH 163
UNCL ASSIFIED \ UNCL ASSIFIED 2% UNCL ASSIFIED
| Y-24920 . Y-24921 3 Y-23964

      
  

Ll
o

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i
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50 Hr. at 1500°F. DPH 174 100 Hr. at 1500°F. DPH 157 1000 Hr. at 1500°F. DPH 148

  

Figure 17. Alloy VI-54, 10 Mo - 10 Cr - 10 Fe - 0.5 AL - 0.5 Mn - 0.06 C - Balance Ni.
Solution-annealed 1/2 hr. at 2100°F, aged at 1500°F. Etchant: Glycera Regia. 500X.

 
-39-

precipitates at the grain and twin boundaries. On the other hand, the
microstructure of the gage length of a creep-rupture specimen of this
alloy illustrated in Figure 18 showed bands of a fine matrix precipi-
tate but still of lesser quantity than in the case of alloy VT-43.
Although it is felt that the difference in quantity of precipitated
carbides between the two alloys is related to a larger difference in
their total carbon contents than indicated by the analyses, it is
recognized that a change in carbon solubility may have been effected by
the variation in chramium and iron contents.

Because of inconsistencies in the factors which affect creep-
rupture behavior, it was not possible to satisfactorily rationalize the
strengths of the alloys. A possible comparison could be made between
the strengths of alloys VT-43 and VT-53 where the principal difference
between the alloys was a 5 per cent increase in the chromium content of
the latter over the former with average grain diameters being compara-
ble. The slightly-higher average strength of alloy VI-53 up to 5 per
cent strain could have been attributed to the increased chramium con-
tent. A more realistic evaluation of the alloys of this group, however,
was the generalization that their strengths were not grossly different
as a result of a summation of the factors affecting creep-rupture
behavior. Their strengths, also, exceeded that of Inconel but were
inferior to that of Hastelloy B which will be shown to be true for all

alloys of series 1.
Lo

o W e e el L P URELASSIFIED
o O e T T N Ry e s Y4B
. W i
. - - g, - R
- ' " & ;
S
" Ll
a 4
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Figure 18. Alloy VI-54, 10 Mo - 10 Cr - 10 Fe - 0.5 Al -
0.5 Mn - 0.06 C - Balance Ni. ILongitudinal section in gage length of
creep~-rupture specimen. Rupture life: 210.9 hr. Etchant: Glycera
Regia. 500X.
 

I Iy

15 Per Cent Molybdenum Alloys with O.5 Per Cent Aluminum - 0.5 Per Cent
Manganese - 0.06 Per Cent Carbon and Varying Percentages of Chromium
and Iron

A summary of the creep-rupture data obtained on the 15 per cent
molybdenum alloys of series I is presented in Table IV and illustrated
by the bar-graph in Figure 19. These data have been presented in the
same manner as the data for the 10 per cent molybdenum alloys discussed
previously. Again, it was found that several factors in addition to
solid-solution effects of molybdenum could have influenced the measured
creep rate, specifically: grain-size variations, dispersed M6C-type
carbides, slight aging tendencies, and variations in the chromium and
iron contents.

Referring to Figure 10, illustrating a solution-annealed micro-
structure of a 15 per cent molybdenum alloy with 0.06 per cent carbon,
it can be seen that a considerable number of M6C-type carbides remalned
undissolved after the solution-annealing treatment at 2100°F. These
particles were very effective in reducing the as-annealed gralin size of
these alloys relative to the as-annealed grain size of the 10 per cent
molybdenum alloys. Average grain-diameter measurements on an unstressed
portion of the creep-rupture specimens shown in Table IV also indicated
a variation in grain size between the individual alloys of this group.
The decrease in average grain diameter in the materials upon increasing
the iron content from 4 to 10 per cent at the three chromium levels has
been attributed to an increase in quantity of M6C-type carbides dis-

persed in the alloys. The widest grain-size variation occurred between

 
TABLE IV

 

SUMMARY OF CREEP-RUPTURE DATA ON Ni-BASE ALLOYS CONTAINING 15 Mo - 0.5 Al -

0.5 Mn - 0.06 C AND VARYING PERCENTAGES OF Cr AND Fe

Test Conditions

 

Temp :
Stress:
Atmos:

1500°F
10,000
Argon

psi

 

 

 

 

Average
Camposition (a) Rupture Elong- Grain
Test Alloy Variable (Wt %) Time to Specified Strain‘>/(Hr.) Life ation Dia.
No No. Cr Fe 1% 2%, 5% 10% (BEr.) (%) (mm )
10-8-3 VT-45 5 L 25,5 L5 ol 150 192.4 16.07 -
10-8-10 30 52 105 173 218.2 16.96 0.0357
10-8-4 VT-L6 5 10 20 35 71 110 152.4 24,10 -
10-8-9 20 36 76 127 200.6 25.00 0.0297
10-8-7 VT-49 7 b 26.5 7.5 97 158 181.8 14.28 0.0400
10-8-25 26 LL 8l 132 174.2 17.85 -
10-8-23 VT-50 7 10 16 29 62 110 225.9 38.37 0.0261
10-8-30 15 27 59 95 115.0 21.42 -
10-8-18 VI-55 10 L 30 59 140 210 259.2 14.28 0.0397
10-8-32 23 Ly oL 148 186.9 18.75 -
10-8-13 VT-56 10 10 24 .5 45.5 102 180 345.1 33.04 -
10-8-16 22 Lo 96 175 LL7.7 46,82 0.0348

 

(a) Microscope readings.
UNCL ASSIFIED

1000 ORNL-LR-DWG. 26325
NOMINAL COMPOSITION (wt %) — — — RUPTURE: 3750 hr
800 BASE: 15 Mo—0.5 Al— 05 Mn— 0.06 C — BALANCE Ni ————— ELONG. 2%
. 5 Cr T Cr 7 Cr {O Cr {0 Cr
600 VARIABLE: o 10 Fe 4 Fe 10 Fe 4 Fe 10 Fe
TEST CONDITIONS
400 TEMP: 150C°F

 

STRESS: 10,000 psi
ATMOS: ARGON —

200

100
80

60

40

20

TIME TO SPECIFIED STRAIN (hr)

 

1

ALLOY: INCONEL VT-45 VT-46 VT-49 VT-50 VT-55 VT-56 HASTELLOY B

Figure 19. Bar-graph of creep-rupture test results obtained at 1500°F, 10,000 psi.,
on Ni-base alloys with 15 Mo - 0.5 Al - 0.5 Mn - 0.06 C and varying percentages of Cr and Fe.

tH

 
Lo

the 7 per cent chromium alloys which are campared in Figures 20 and 21.
In general, the difference in guantity of dispersed.MéC—type carbides
was supported by the analyzed carbon contents of the alloys.

Comparison of the range and average hardnesses of these alloys
in the annealed and aged conditions is shown in Table V, and the
average hardnesses are plotted in Figure 22. Compared to the 10 per
cent molybdenum alloys, the 15 per cent molybdenum alloys were harder
and showed less aging response. The 15 per cent molybdenum alloys
contained a greater number of M6C—type carbides which did not dissolve
upon solution-annealing. As a result, the amount of total carbon
available for reprecipitation as carbide particles was less upon aging
at 1500°F.

A correlatlion between microstructure and hardness was found for
all alloys of this group. Because similar microstructures were devel-
oped within the materials upon aging and creep-rupture testing,
representative photamicrographs are presented only for the materials
with the lowest and highest alloy contents. The microstructures of the
aged material with the lowest alloy content, VI-45 (15 Mo - 5 Cr -

4 Fe - 0.5 A1 - 0.5 Mn ~ 0.06 C - balance Ni) are shown in Figure 23.
After five hours at 1500°F, there waé an indication of a grain=boundary
reaction. The only apparent change in microstructure for aging times
greater than 5 hours was for the grain-boundary particles to became
scmewhat more discrete and show coalescence after aging for 1000 hours.
The application of stress did not have a measurable. effect on the

microstructure of this alloy, as shown in Egure 24. The microstructures

 
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L i 2 3 = 3 R ke T S Hn=tdl o ae

Figure 20. Alloy VI-49, 15 Mo - 7 Cr - 4 Fe - 0.5 Al - 0.5 Mn
0.06 C - Balance Ni. Longitudinal section in grip of creep-rupture
specimen. PRupture life: 181.8 hr. FEtchant: Glycera Regia. 100X.

o e A e bl

i e =

1"4;"‘--_...1 = S i o-hlr
- - sl e, -

Gl i 2y g .__!.-_-,_....Il o e _.-1p:.‘ ; __;...r.- 2, LT

2w

 

Figure 21. Alloy VI-50, 15 Mo - 7T Cr - 10 Fe - 0.5 Al
0.5 Mn - 0.06 C - Balance Ni. Longitudinal section in grip of creep-

rupture specimen. Rupture life: 225.9 hr. Etchant: Glycera Regla.
100X.

 
DIAMOND PYRAMID HARDNESS DATA ON Ni-BASE ALLOYS CONTAINING 15 Mo -

TABLE V

 

0.5 AL - 0.5 Mn - 0.06 C AND VARYING PERCENTAGES OF Cr AND Fe.
SOLUTION-ANNEALED 1/2 HR. AT 2100°F, AGED AT 1500°F.

 

 

 

 

 

Comoatiion olutton hee s 1500°8

Alloy (Wt %) at 2100°F 5 Hr. 25 Hr. 50 Hr. 100 Hr. 1000 Hr.

No. Cr Fe  Range Avg. Range Avg. Range Avg. Range Avg. Range Avg. Range Avg.
Vr-45 5 4L 173177 174 187189 188 1834189 187 181-188 185 176-181 179 169188 180
vT-46 5 10 176191 184 196-201 199 192-194 193 192-194 193 193-197 195 189-194 191
yr-49 7 L 173-181 178 202205 203 197-203 199 189-196 192 192-198 195 180-188 164
VI-50 7 10  182-198 191 198199 198 187-192 190 198-198 198 187-188 187 186-201 194
VI-55 10 L 181196 188 20106 204 202-203 203 206-207 207 199-205 202 189197 193
VI-56 10 10 185-189 187 196—205 201 201—-206 203 198-201 199 201202 201 188-196 192

 

—91-{_
UNCL ASSIFIED

ORNL-LR-DWG, 26321

 

NOMINAL COMPOSITION {wt %)
230 BASE: {5 Mo—0.5 At—0.5 Mn—0.06 C—BALANCE Ni
5Cr 5Cr 7 Cr 7 Cr {0 Cr 10 Cr
VARIABLE : | 4 Fe 10 Fe 4 Fe 10 Fe 4 Fe 10 Fe B
220
o 2i0
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< 0| < Q < o < o < o a o
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160
ALLOY: VT -45 VT -46 VT ~49 VT-50 VT-55 VT-56

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

LY

Bar-graph of average hardness of Ni-base alloys with 15 Mo - 0.5 Al -

Figure 22.
Solution-annealed 1/2 hr, at 2100°F,

0.5 Mn -~ 0.06 C and varying percentages of Cr and Fe.
aged for 5, 25, 50, 100, and 1000 .hr. at 1500°F.
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/
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Solution Annealed.

F o™

 

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DPH 174

   

5 Hr.

UNCL ASSIFIED

Y-2483
0% $
.

DPH 1868

at 1500°F.

2 UNCLASSIFIED h S UNCL ASSIFIED
; o Y-24834 Y-24833
- o ey = S
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o e = =i .. o '.. . o - -'. g™ @ o ey e :;b S
S Tan, Ny s E . e S W y
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Pq'. ? TH o :go l"fl & '-a :‘n.":.* . A i y
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a - ’.1::' °"‘-nd . - n' ¥ - _‘,3
‘_‘ “L - ‘ ot fi 'l
f L 4 \ a- - % -
- - Sy,

50 Hr. at 1500°F.

Figure 23.

!
v

DPH 185

100 Hr. at 1500°F.

Alloy VI-45, 15 M
Solution-annealed 1/2 hr. at 2100°F, aged at 1500°F.

DPH 179

UNCL ASSIFIED

2 Y-24835
pe\ - e “ L = ansaa :-""'\-:""-"h
r
a - & " - X -
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5 - B o
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L *
Pt st
25 Hr. at 1500°F. DPH 187
BTEY T . 3T R 9 UNCLASSIFIED
. 50 S‘W@wfififiw
] o y
o= & &P !} Wfi v = ‘: i
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.‘ -~ A.:._Jf - - '. Q\
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-, }"“'a sn 4.9 o Gt _wud"':“.‘"
Y . P ST
g i

1000 Hr. at 1500°F. DPH 180

0o-5Cr -4 Pe - 0.5 AL - 0,5Mn -~ 0.06 C - Balance Ni.
Etchant :

Glycera Regla. 500X.

 
F T PR B e IR 18  UNCLASSIFIED
: S ® . S AIhie d Y-M'{S
e i(" - Bdg & /v 4
ol Y P — e, . =
a, - h ¥ i \
3 J“_ ‘ '! < b"fi . i N a £ uE ;
p- g s x ;-,‘ :_':-u " :\ Q. ‘e G ¥ s '_'. ._' ;
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i " i' p :_,;_‘.“l ';‘ 3 -\' = .‘-" '.l‘EP ® "‘
¢ it e . - & . e
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: % = I PV " ML -
.~ e et .- - - c - 0w L [\f’h
i LY '{L" " .
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g © . -F_ Y ot oy, / ,: o B 2 .-b
-

Figure
0.5 Mn - 0.06
creep-rupture
Regia. 500X.

24k, Alloy VI-45, 15 Mo - 5 Cr - 4 Fe - 0.5 Al -
C - Balance Ni. Longitudinal section in gage length of
specimen. Rupture life: 218.2 hr. Etchant: Glycera
 

-50-

developed by the material with the highest alloy content, VI-56 (15 Mo -
10 Cr - 10 Fe - 0,5 AL -~ 0.5 Mn - 0.06 C - balance Ni), after aging and
creep~-rupture testing are shown in Figures 25 and 26, respectively.

The creep strength of the 15 per cent molybdenum alloys was
greagter than that of the 10 per cent molybdenum alloys. The principal
contributing factors to this increased strength were undoubtedly the
solid-solution strengthening effect of molybdenum, and the carbide
dispersion strengthening effect. These were of sufficient magnitude to
overcome the advantage of coarser grain size held by the 10 per cent
molybdenum alloys.

Increasing the chromium and iron contents of the 15 per cent
molybdenum alloys did not bring about pronounced strengthening effects.
In fact, there was a slight weakening of the alloys upon increasing the
iron content from 4 to 10 per cent at a constant chromium concentration.
An explanation based upon the finer grain size of the high-iron alloys
does not follow since the cause of the finer grain size was attributed
to an increase in quantity of M6C-type carbides in the alloys; the
latter would be expected to over-ride the influence of grain size upon
the creep~rupture behavior. The exact cause of the strength differences

noted is not clear, therefore, fraom the data which have been obtained.

 
| - HUNCI;'AESEz[f_El.ED R \‘r;, cw e
e , - % g - - 24837
P : : %‘;&v" e £ Bpn. s s
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i

Solution Annealed. DPH 187 5 Hr. at 1500°F. DPH 201
o AR g B s el e s UNCLASSIFTED 2l UNCL ASSIFIED
L s ifoed Y-24839 o o ;g Y-24840
o o e . S T a e * o . )y e . s ! i
o o A e, P G ‘lg;%? :P. R& %'Oop%rfl.
= X . = N5 ¥oa
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oo g TP - 'Q : ; o i .‘ ' o
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# -y a5 "'n ’u - G'F. . & & s dfl*
g v Y 4, & 02 COOS S [ -
a = o 6 ] a " s '0 o .D-q," -
- oG = el a3k g'2" ° o = g
: " s
%

50 Hr. at 1500°F. DPH 199 100 Hr. at 1500°F. DPH 201

Figure 25. Alloy VT-56, 15Mo~100r-10Fe-o5A1
Solution-annealed 1/2 hr. at 2100°F, aged at 1500°F. Etchant:

po- 0O 3 o o LINCL ASSIFIED
P e ” . Y&afi_BB
o Ty ©
TR . ,om Q‘Chg.-n-dfl-
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oo e et 5 TR T -
oy "fl‘-‘!. r"’):". - %h L On
= -
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\ : _UNCL ASSIFIED
& e Y.24001
o e - - ¥ np
okl " v . .
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- 0.5 Mn - 0.06 C - Balance Ni.
Glycera Regia. 500X.

TS
52

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3 ..E.r‘ ‘: .%.. o "E ',u'e 3 o ‘:'- . . Bups
& - I ‘4 -t Chatls x w " / (-' e
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o . ot .‘C-_. o - £ ) .
" - s o -
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a 5 ;%H- s Dredt TS -\ » LR '{,3’ *\-fif"‘
£ o R - -
3 oA TN
Ry B e o }’\\ o
-
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- -3 e c mZpg _- bi::g; — \-\q.{' o e

Figure 26. Alloy VT-56, 15 Mo - 10 Cr - 10 Fe - 0.5 Al -
0.5 Mn - 0.06 C - BRalance Ni. Longitudinal section in gage length of
creep-rupture specimen. Rupture life: UL47.7 hr. Etchant: Glycera
Regia. 500X.

 
 

“53-

20 Per Cent Molybdenum Alloys with O.5 Per Cent Aluminum - 0.5 Per
Cent Manganese - 0.06 Per Cent Carbon and Varying Percentages of

Chromium andrIron

 

Hot-rolling of the 20 per cent moliybdenum alloys was less
successful than for the lower molybdenuri~-content alloys. As a result,
creep-rupture specimens could be prepared from only three of the four
alloys listed in Table VI where a summary of the creep-rupture results
is presented. These data are plotted in bar-graph form in Figure 27.
As in the case of the 10 and 15 per cent molybdenum alioys of this
series, an attempt was made to interpret the data in terms of the
variables which could be influencing creep-rupture behavior,

The solution-annesled grain size of this group of alloys was
smaller than that of the 10 and 15 per cent molybdenum alloys and was
attributed to the increase in the quantity of M6Cutyp carblides present
in the microstructure. A very slight variation in grain size between
the individual alloys was alsc noted as shown in Table VI.

Hardness and metallographic stucies were conducted on all four
alloys of this group after aging at 1500°F. Table VII shows the range
and average hardness of the various coupons of each alloy after aging;
the average hardnesses are shown graphically in Figure 28,

The aging response of these alloys showed evidence that a phase
boundary had been crossed by the 10 per cent iron alloys at both
chromium concentrations. It should be mentioned, however, that both
10 per cent iron alloys contained 7 per cent chramium according to the

reported analyses. The precipitation of the additionsl phase in these

 
TABLE VI

 

SUMMARY OF CREEP-RUPTURE DATA ON Ni-BASE ALLOYS CONTAINING 20 Mo - 0.5 AL -

0.5 Mn - 0.06 C AND VARYING PERCENTAGES OF Cr AND Fe
‘ Test Conditions

1500°F
10,000 psi
Argon

 

Temp :

Stress:

Atmos:

 

 

 

 

Time to Specified Average
Camposition (a) Rupture Elong- Grain
Test Alloy Variable (Wt %) Strain‘~/(Hr.) Iife ation Dia.
No. Ne. Cr Fe " 1% 2% 5% 10% (Hr.) (%) (mm)
10~-8-14 V=57 5 b ol L7 105 180 284 .2 21.43 -
10-8-26 33.5 60 122 205 293.1 22,31 0.0318
10-8-24 VT-58 5 10 26 53 130 260 660.5 h7.31 0.0247
10-8-27 18 38 100 200 hhg, 1, 39.28 -
10-8-22 VT-59 T b 30 54 118 200 34k, 38.39 0.0289
10-8-28 25,5 b5 95 162 298.0 51.78 -
VT-60 7 10 (No tests)

 

(a) Microscope readings.

-'I-(Q-
55

UNCLASSIFIED

-LR-DWG, 26326
2000 ORNL

NOMINAL COMPOSITION {(wt %)
BASE: 20 Mo— 0.5 Al—=0.5 Mn—0.06 C —BALANCE Ni

 

. S Cr 5 Cr Cr
1000 VARIABLE: 4 Fe 10 Fe 4 Fe
800 _ RUPTURE: 3750 hr
TEST CONDITIONS ELONG. 12%
600 TEMP: 1500°F

STRESS: 10,000 psi
ATMOS: ARGON

400

200

100
80

60

40

20

TIME TO SPECIFIED STRAIN (hr)

 

ALLOY: INCONEL vT-57 VT—-358 VT-59 HASTELLOY B

Figure 27. Bar-graph of creep-rupture test results obtained at
1500°F, 10,000 psi., on Ni-base alloys with 20 Mo - G.5 Al - 0.5 Mn -
0.06 C and varying percentages of Cr and Fe.
 

TABLE VITI

DIAMOND PYRAMID HARDNESS DATA ON Ni-BASE ALLOYS CONTAINING 20 Mo -
0.5 Al - 0.5 Mn - 0.06 C AND VARYING PERCENTAGES OF Cr AND Fe.
SOLUTION.ANNEALED 1/2 HR. AT 2100°F, AGED AT 1500°F. |

 

Composition Solution Aged at 1500°F

 

 

 

Variable Annealed
Mloy (Wt %) at 2100°F 5 Hr., 25 Hr. 50 Hr. 100 Hr. 1000 Hr.
No. Cr Fe Renge  Avg. Range Avg. Range Avg. Range Avg. Range Avg. Range  Avg.
VI-57 5 L 201206 203 205215 212 206-215 212 203206 205 202205 204 198210 20L
VI-58 5 10 198205 201 207-210 209 202210 206 207210 209 215-218 216 235~2L40 238
VIT-59 7 L 201-215 209 207207 207 215922 217 210219 216 215~p22 218 207-215 212
VI-60 7 10 194-202 198 207-215 211 201210 207 212-218 214k 207-219 213 235-236 236

 

_96_
27

UNCL ASSIFIED
ORNL«L R«DWG. 26322

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

260
NOMINAL COMPOSITION (wt %)
BASE: 20 Mo—0.5 AI—0.5Mn—0.06 C—BALANCE Ni o
i 5Cr 5Cr 7 Cr 7 Cr |
VARIABLE:| 2 £ 10 Fe 4 Fe 10 Fe
240
" _
N —
W L
2
O
&
T 230 L
o
=
q —
o
).
a
o 220 i
Z r———
O —
E — e
« — -
o ||
210 ] - ] -
200 oo | ol oe . Ll ]
<) o Ll o o (1= e Lo — < o | o
%mm-‘:ff %tflm-‘:-CE %LOLO'C.C_C ,I-OL-::')'C_C.C
a| |[NO|plo | |NOlolo | |NClojo Z| (NIO|polo|
Violo OO Olold Z Oio|lo
2 TS Tiol € <|o
190
ALLOY: VT—57 VT—58 VT—59 VT—60

Figure 28. Bar-graph of average hardness of Ni-base alloys with
20 Mo - 0.5 AL - 0.5 Mn - 0.06 C and varying percentages of Cr and Fe.
Solution-annealed 1/2 hr. at 2100°F, aged for 5, 25, 50, 100, and 100COC
hr. at 1500°F,

 
-58-

alloys brought about a steady increase in hardness with aging time at
1500°F which could be correlated with an increase in grain~boundary
pregipitate, The hardness data for the low-~lton alloys at the two
chromium contents agreed well with the trend previously noted, in that
the response to carbide aging was only slight due to the tie~up of
carbon in the M6C-type carbides which did not dissolive during the
solution-annealing *treatment., Consequently, the difference between the
solution-annealed hardness of these two alloys and thelr peak hardness
upon aging was less than that noted for the 15 per cefit molybdenum
alloys. The general hardness level of the 20 per cent molybdenum alloys
was greater than for the lower molybdenum alloys.

The course of aging of alloy VI-59 (20 Mo - 7 Cr - 4 Fe -
0.5 AL - 0.5 Mn - 0.06 C - balance Ni) at 1500°F is illustrated in
Figure 29. A grain-boundary reaction had occurred after five hours,
and discrete grain-boundary particles were observed after 25 to 1000
hours, Fligure 320 shows there was little effect of stress on the micro-
structure of this alloy. The solution-annealed and aged microstructures
developed by alloy VI-60 (20 Mo = 7 Cr - 10 Fe - 0.5 AL - 0,5 Mn -
0,06 C - balance Ni) are illustrated in Figure 31L. Precipitation of an
additional grain-boundary phase was observed with Increasing aging time.
Figure 32 more clearly shows this phase as it developed in the gage
length of a ereep-rupture specimen of alloy VI-58 (20 Mo - 5 Cr -
10 Fe - 0.5 AL - 0.5 Mn - 0.06 C - balance Ni).

If the average time, i.e., average of the two tests per alloy,

to 1 per cent strain was used as a criterion for determining the

 

 
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polution Annealed. DPH 209 5 Hr. at 1500°F. DPH 207 25 Hr. at 1500°F. DPH 217
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50 Hr. at 1500°F. DPH 216 100 Hr. at 1500°F. DPH 218 1000 Hr. at 1500°F. DPH 212

Figure 29. Alloy VI-59, 20 Mo - 7T Cr - 4 Fe - 0.5 Al - 0.5 Mn - Q.06 C - Balance Ni.
Solution-annealed 1/2 hr. at 2100°F, aged at 1500°F. Etchant: Glycera Regia. 500X.
60

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Figure 30.

0.5 Mn - 0.06 C - Balance Ni,.
creep-rupture specimen.

Regia. 500X.

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Alloy VI-59, 20 Mo - 7T Cr - 4 Fe - 0,5 Al -
Longitudinal section in gage length of
Rupture life: 344.1 hr. Etchant: Glycera

 
 

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Figure 3L. Alloy VI-60, 20 Mo - 7 Cr - 10 Fe - 0.5 AL - 0.5 Mn - 0.06 C - Balance Ni.
Solution-annealed 1/2 hr. at 2100 °F, aged at 1500°F. BEtchant: Glycera Regia. 500X.

 
62

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B g,

: & £ = i ‘3‘3 D'J.-_‘;l:' PO .:%‘
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Figure 32. Alloy VT-58, 20 Mo - 5 Cr - 10 Fe - 0.5 Al -

0.5 Mn - 0.06 C - Balance Ni. Longitudinal section in gage length of

creep-rupture specimen. Rupture life: 660.5 hr. Etchant:
500X .

Regia.

Glycera

 
-63-

strength of the alloys in which only carbide precipitation was present,
thére was little difference between the 20 per cent molybdenum alloys
and the corresponding 15 per cent molybdenum alloys. However, the
average times to stralns greater than 1 per cent showed the 20 per cent
molybdenum alloys to be stronger.

Comparing the data between the individual 20 per cent molybdenum
alloys, there was little difference in the strengths of alloys VI-59
and VI-57. It was Intended that alloy VI-59 represent a 2 per cent
increase in the naminal chromium content. The analyzed compositions of
these two alloys, however, showed the higher chramium alloy to contain
approximately 3 per cent less molybdenum. The fact that no difference
in the amount of precipitate was noted within the two materials indi-
cated the creep strength to be more dependent upon structure than the
variation in molybdenum and chromium contents.

It was interesting to note the behavior of the 10 per cent iron
alloy, VI-58, in creep~-rupture. Average times for strains up to 1 per
cent were less for this alloy than either of the alloys VI=57 or VI-59,
However, the times required for strains between 2 and 5 per cent were
canparable for all three alloys. For greater amounts of strain the
greater strength of alloy VI-58 was evident by the progressively
longer times for a given strain. The rupture life of alloy VT-58 was
significantly longer., This behavior was probably associated with the
slightly finer grain size of alloy VI-58 compared with that of the
other two alloys coupled with instabilities resulting from the pre-

cipitation of the non-carbide grain-boundary phase. These factors

 
-6l

caused an Iincreased initlal creep rate. The non-carblide grain-boundary

phase later added to the strength in the final stages of test.

"High-Purity" Alloys

Because the carbon intentionally added to the alloys of series I
was very effective in introducing into these materials several of the
factors which affect creep-rupture behavior, strengthening effects
attributable to chramium and/or iron could not be established ffam the
deta with any degree of certainty. Therefore, the alloys of se}ies IT
were prepared with no intentional carbon addition. The alloys were
intended to show the Iinfluence of increasing amounts of chramium upon

the strength of a 15 Mo - balance Ni binary alloy, and the influence of

increasing amounts of iron upon the strength of a 15 Mo - 7 Cr - balance

- N1 ternary alloy.

The creep-rupture data obtained on the alloys of series II are
sumarized in Table VIII. Considering the influence of chramium, it
would be Inferred from the bar-graph of creep-rupture results shown in
Figure 33 that more than 3 per cent chromium in a 15 Mo = balance Ni
binary alloy brought about an increase in strength. Unfortunately, it
was necessary to consider also the effects of grain-size variations and
a precipitation reaction in the alloys containing 5 per cent or more of
chramium.

Grain-size measurements tabulated in Table VIII made on the
creep=rupture specimens of the straight chromium-bearing alloys showed
a sharp decrease in grain size when they contained 5 per cent or more

of chramium. Apparently, the stress-relieving treatment at 1600°F for

 
TABLE VIII

SUMMARY OF CREEP-RUPTURE DATA ON Ni-BASE ALILOYS CONTAINING
15 Mo AND VARYING PERCENTAGES OF Cr AND Fe
Test Conditions

Temp : 1500°F

Stress: 10,000 psi
Atmos ; Argon

 

 

 

 

' Average
Camposition (a) Rupture Elong- Grain
Test Alloy Variable (Wt %) Time to Specified Strain‘?/(Hr.) Life ation Dia.,
No. No. Cr Fe 0.5% 1% 2% 5% 10% (Hr.) (%) (mm)
10-8-42 VI-90  PBase - 2.7 3.9 6.2 10.7 - - 12.6 7.1u" 0.,0952
'10-8-49 4.5 6.6 9.6 15.5 - 19.7 8.03 -
10-8-39 V-89 3 - 3.5 4.9 6.6 - - 7.8 3.57 0.0917
10~-8-50 "3.5 4.9 6.7 9.5 - 10.1 6.25 -
10-8-38 vT-88 5 - T2 10.2 1k4.5 - - 19.0 3.57 -
10-8-40 8.4 11.2 1L4.5 - - 17.4 3.57 0.0492
10-8-36 yr-87 7 - 9 12.1 16 20.7 ~ 20.8 5.35 -
10-8-41, 9.6 13.3 19 29.5 - 31.8 . 6.25 0.0591
10-8-43 VT-86 10 - - 15,5 27 56 95 122.6 16.96 -
10-8-46 11.7 18.8 31.5 68 115 159.9 18.75 0.0510
10-8-L7 VT-91 7 L 7.5 10.5 13.8 18.2 - 18.4 5.35 0.0869
10-8-48 8.8 12.3 17.5 ol .7 - 25.6 7.1k -

ngg-

 
 

TABLE VIII (continued)

SUMMARY OF CREEP-RUPTURE DATA ON Ni-BASE ALLOYS CONTAINING
15 Mo AND VARYING PERCENTAGES OF Cr AND Fe

Test Conditions

Temp : 1500°F
Stress: 10,000 psi
Atmos: Argon

 

 

 

 

Average
Composition (a) Rupture Elong- Grain
Test Alloy Variable (Wt jfl_ Time to Specified Strain a (Hr.) Iife ation Dia.
No. No. Cr Fe 0.5% 1% 2% 5  10% (Ar.) (%) (rom )
10-8-Lk VT=-92 7 T 13.2 19 29 50 67 68,1 10.71 0.0641
10-8-34 1k 20.8 32 57 80 80.3 10.71 -
10-8-45 VT-93 7 10 12.5 19.2 30 5k - 75.8 9.82 0.0641
10-8-35 12.5 19,2 30 5k - 73.7 8.92 -

 

(a) Microscope readings.

..-99..
TIME TO SPECIFIED STRAIN {hr)

1000
800

600

400

VARIABLE:

200

100

@
O

D
O

H
O

Ny
o

ALLOY: VT-90

Figure 33.

 

67

UNCL ASSIFIED
ORNL-LR-DWG, 26327

NOMINAL COMPOSITION (wt %)

BASE: 15 Mo—BALANCE Ni

5Cr 7 Cr

TEST CONDITIONS

TEMP: (500°F
STRESS: 10,000 psi
ATMOS: ARGON

VT-89 VT-88 VT-87 VT-86

Bar-graph of creep-rupture test results obtained

at 1500°F, 10,000 psi., on Ni-base alloys with 15 Mo and varying

percentages of Cr.

 
~-68-

one-half hour given all the materials prior tc machining creep-rupture
specimens precipiteted the grain=boundary phase which, until effec-
tively dissolved, retarded grain growth in the specimens during the
solution-annealing treatment.

The range and average hardness of the chromium-bearing ternary
alloys after a solution-annealing treatment as well as after aging
100 hours at 1500°F are shown in Table IX. The average hardness data
are presented in bar-graph form in Figure 34. Increasing the chromium
content caused an increase in hardness of the ternary alloys, although
it should be remembered that the hardness measurements included a
graine-size variable as well as a camposition variable. The slight
increase in hardness as a result of the 1500°F aging treatment in the
alloys containing 5 per cent or more of chromium could be correlated
with the precipitation which occurred in these alloys. |

Figures 35, 36, and 37 show the microstructure of alloy VT-90

(15 Mo - balance Ni) after having been solution-ennealed, aged 100 hours

at 1500°F, and creep-rupture tested at 1500°F. 1In all cases the struc-

ture is that of a solid-solution alloy. These microstructures were also

representative of alloy VI-89 (15 Mo - 3 Cr - balance Ni) under the same

conditions.

Figures 38, 39, and 40 illustrate the microstructures found in
alloy VI-87 (15 Mo - 7 Cr - balance Ni) after solution-annealing, aging,
and creep-rupture testing. This series of photamicrographs depicts the

grain-boundary precipitate typical of that found in the alloys.

 
-69-

TABLE IX

DIAMOND PYRAMID HARDNESS DATA ON Ni-BASE ALIOYS CONTAINING 15 Mo

AND VARYING PERCENTAGES CF Cr AND Fe.

SOLUTTON-~ANNEALED

1/2 HR. AT 2100°F, AGED FOR 100 HR. AT 1500°F.

 

 

 

 

 

 

Camposgition Solution Annealed - Aged 100 EHr,

Alloy Variable (Wt %) __at 2100°F at 1500°F

- No., Cr Fe Range AvE. Range AvVg.
VT'-=90 Base 139~141 140 132-133 132
vT-89 3 - ik L6 145 141145 143
vT-88 > = 155155 155 156159 157
vT-87 7 - 151159 155 157160 158
vT-86 10 - 159163 161 163-166 165
VT-91 T 4 148-151 150 151-156 15k
VI-92 7 T 151157 155 155-162 159
VI-93 T 10 148~155 152 153-158 155

 

 
NOMINAL COMPOSITION (wt %)
BASE: 15 Mo—BALANCE Ni
180
7 Cr 7 Cr 7 Cr
VARIABLE: {BASE 3 Cr 5 Cr 7 Cr 10 Cr 4 Fe 7 Fe 10 Fe
170
W
D
Ll
=z —
o 160
a s
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£ Zlo Z|2
Zlo 2O <
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120

ALLOY: VT-90 VT-89 VT-88 VT-—-87

UNCL ASSIFIED

ORNL-L R-DWG, 26323

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

 

VT—-86 VT-—91

VT—92 VT-93

Figure 34. Bar-graph of average hardness of Ni-base alloys with 15 Mo
Solution-annealed 1/2 hr. at 2100°F,

and varying percentages of Cr and Fe.
aged for 100 hr. at 1500°F.

oL
" UNCL ASSIFIED
Y-24005

// ey

Figure 35. Alloy VI-90, 15 Mo - Balance Ni. Solution-annesled
1/2 hr. at 2100°F. DPH 140. Etchant: Glycera Regia. 500X.

UNCLASSIFI ED
Y-24931

Figure 36. Alloy VI-90, 15 Mo - Balance Ni. Solution-annealed
1/2 hr. at 2100°F, aged for 100 hr. at 1500°F. DPH 132. Etchant:
Glycera Regia. 500X.

 
UNCL ASSIFI ED

Pl

_. /

R

Figure 37. Alloy VI-90, 15 Mo - Balance Ni. ZLongitudinal
section in gage length of creep-rupture specimen. Rupture life:
12.6 hr. Etchant: Glycera Regia. 500X.

 
 

Figure 38.

Alloy VT-87, 15 Mo - 7 Cr - Balance Ni.
annealed 1/2 hr. at 2100°F. DPH 155.

Solution-
Etchant: Glycera Regia. 500X.

UNCL ASSIFIED
Y-24930

\\‘,rl

Figure 39. Alloy VI-87, 15 Mo - 7 Cr - Balance Ni. Solution-
annealed 1/2 hr. at 2100°F, aged for 100 hr. at 1500°F. DPH 158.
Etchant: Glycera Regia. 500X.

 
T

. : UNCLASSI FIED
€ N~ NS08

. : d_f'. -
{
JARNT S v
'-fl "'. W .
y 3

}.J‘
\

Figure 40. Alloy VT-87, 15 Mo - 7 Cr - Balance Ni. Longi-
tudinal section in gage length of creep-rupture specimen. Rupture
life: 31.8 hr. Etchant: Glycera Regia. 500X.
 

_75..

containing 5, 7, and 10 per cent chramium as a result of aging or
creep-rupture testing at 1500°F.

Because of the inconsistencies in the grain sizes and the amount
of grain-boundary precipitate in this group of alloys, a comparison of
their relative strengths could at best be made between the O and 3 per
cent chromium alloys, and between the 5, 7, and 10 per cent chromium
alloys. With this scheme of evaluation, it was concluded that a 3 per
cent chromium addition added nothing to the strength of the binary base
composition. Similarly, only a slight increase in strength was realized
by increasing the chromium content from 5 to 7 per cent; but an increase
to 10 per cent chromium showed a more significant strengthening effect,
particularly for the time required to reach a strain of 2 per cent or
more .

The creep-rupture data obtained on the alloys prepared to show
the influence of increasing amouhts of iron upon the strength of a
15 Mo = 7 Cr - balance Ni ternary composition are plotted in Figure Ll.
Interpretation of the data was again complicated by the necessity of
considering grain size and precipitation variables in addition to
camposition.

Grain-size measurements made on a creep=rupture spécimen of each
of the iron-bearing alloys are tabulated in Table VIII. The wvariaticn
noted has been attributed to the difference in the amount of precipi-
tation which:occurred in these alloys during the stress-relieving
treatment prior to machining the creep-rupture specimens. It will be

shown that the volume of grain-boundary particles precipltated at 1500°F

 
UNCL ASSIFIED

ORNL -L R-DWG. 26328
1000 R

800

600 NOMINAL COMPOSITION (wt %)
BASE: 15 Mo—7 Cr —BALANCE Ni

400

VARIABLE: | BASE

200

TEST CONDITIONS

TEMP: 1500°F
100 STRESS: 10,000 psi
o ATMOS: ARGON

 

(hr)

60

40

20

TIME TO SPECIFIED STRAIN

10

 

ALLOY: VT-87 VT-91 VT-92 VT—-93

Figure 41l. Bar-graph of creep-rupture test results obtained at
1500°F, 10,000 psi., on Ni-base alloys with 15 Mo - 7 Cr and varying
percentages of Fe.

 
-T7 -

in the alloy of lowest 1ron content was less than that found in the
T and 10 per cent iron alloys.

The range and average hardness of the alloys after solution -
annealing at 2100°F and after aging 100 hours at 1500°F are shown in
Table IX. The average hardness data have.been plotted in bar-graph
form in Figure 34. There was no significant trend in hardness as a
function of increasing iron content; the slight rise in hardness of
each alloy after aging 100 hours at 1500°F has been attributed to the
grain-boundary precipitation which took place in the alloys.

Microstructures developed by alloy VT-91 (15 Mo - 7 Cr - 4 Fe -
balancé Ni, chemical analysis indicated 5 Cr and 2 Fe) after having been
solution-annealed, aged 100 hours af 1500°F, and creep~-rupture tested at
1500°F are shown in Figures 42, 43, and 44. Although a small quantity
-of grain-boundary precipitate was observed in the aged coupon, virtu-
ally none was found in the creep-rupture specimen.

The T and. 10 per cent iron alloys precipitated a larger amount
of the grain-boundary phase which was observed in both the 100-hour
aged specimens and the creep-rupture specimens. This is evident when
Figures 45, 46, and 47, microstructures of alloy VI-92 (15 Mo - 7 Cr -
10 Fe - belance Ni) are ccmpared with the previous three figures. The
amount of grain-boundary precipltate observed in these alloys could be
correlated with their carbon analyses. Both the 7 and 10 per cent iron
alloys contained approximately twice the amount of carbon as the 4 per

- ¢cent iron alloy.

 
UNCL ASSIFIED

! Y-25105
~—
= N

Alloy VI-9l, 15 Mo - 7 Cr - 4 Fe - Balance Ni.

Figure 42.
Etchant: Glycera

Solution-annealed lf? hr. at 2100°F. DPH 150.
Regia. 500X.

7 UNCL ASSIFIED
\"‘\.H T*@D 8
N -
P g
i i
-
g r
i
R el : 7
e | e -
«
o /
. /
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R f
] i = 1'{ f
- - |
!
= fif- by
- \

Figure 43. Alloy VI-9l, 15 Mo - 7 Cr - 4 Pe - Balance Ni.
Solution-annealed 1/2 hr., at 2100°F, aged for 100 hr. at 1500°F.
DPH 154. Etchant: Glycera Regia. 500X.

 
UNCL ASSIFIED

Y-24928
'
|
b
'
/ ™ !
¢ \h ‘
Jf - —
/ - o,
,-‘h - Mo
S
Z__r:—‘—‘_'_‘
' {
s

Figure 4k. Alloy VT-91, 15 Mo - 7 Cr - 4 Fe - Balance Ni.
Longitudinal section in gage length of creep-rupture specimen.
Rupture life: 18.L4 hr. Etchant: Glycera Regia. 500X.

 
80

\ UNCL ASSIFIED
‘\\ . e Y-24011
)

—

Figure 45. Alloy VI-93, 15 Mo - 7 Cr - 10 Fe - Balance Ni.
Solution-annealed 1/2 hr. at 2100°F. DPH 152. Etchant: Glycera
Regia. 500X.

UNCL ASSIFIED
?-24923

iy e

= | :
T . \\ 7 ‘\‘
¢ ',g_-'."' . ;
f / S ;
~ . A N\ " e
N / /
\ /:" F .,
f’ X I 1
* /
v |
'-b“ /' |l'
(
XS , ;
\ 2 o~ . e - ‘.‘
’ , Vi \
P ‘1. ~ g \
J i e . ; E
# ; = 8 " \ \' EAL =
J/ . € ../ ‘I-

Figure 46. Alloy VI-93, 15 Mo - 7 Cr - 10 Fe - Balance Ni.
Solution-annealed 1/2 hr. at 2100°F, aged for 100 hr. at 1500°F.
DPH 155. Etchant: Glycera Regia. 500X.

 
» . . UMCL ASSIFIED
. Y-24927
i
I" » . ®
- "‘
A
. .
- "‘ i, TR
LY . f
. !.’ .\-
\ o /
. , - *e

Figure L47. Alloy VI-93, 15 Mo - 7 Cr - 10 Fe - Balance Ni.
Longitudinal section in gage length of creep-rupture specimen.
Rupture life: 75.0 hr. Etchant:; Glycera Regia. 500X.

 
The presence of other variapbles in addition to campesition
caused difficulty in attempting to determine the influence of iron upon
the strength of the alloys. For example, comparing the strength of the
ternary base camposition, VI-87, with that of the 4 per cent iron allcy,
VI-91, it was found that both materials were comparable in strength.
However, such a compariscn necessitated an accounting not only for the
difference in camposition, but also for the differences in grain size
and amcunt of precipitation present in the materials. It was, then, a
summation of the variables affecting creep-rupture behavior which
equalized the strength of the alloys.

It was possiblie to make a direct comparison of the effect of
increasing iron content between the allcoys containing 7 and 10 per cent
iron. The fact that ro increase in strength was realized by the further
increase of 3 per cent iron showed this element to have an insignifi-

cant effect as a solid-solution strengthener.

DECARBURIZATION STUDIES

A caomparison of the microstructures of the 15 per cent molyb-
denum alloys with and without an irtentional carbon addition left little
doubt that carbides accounted for the bulk of second phase materlal
observed in these alloys, since the quantity present was related to the
carbon content. Also, the fact that the nickel-molybdenum-chromium
ternary compositions of series IT were located well within the alpha

phase boundaries of the 1508°F section of this gystem proposed by

 
-83-

Lundy and Stansbury7 indicated that the precipitates found at 1500°F in

these alloys were carbides. On the other hand, the observed increase
in the amount of precipitated particles with increasing percentages of
iron added to the 15 per cent molybdenum - 7 per cent chromium -
balance nickel alloy warranted more conciusive phase identification in
the iron-bearing alloys of series II. It has already been shown by
creep-rupture tests, hardness studies, and metallographic means that
precipitation of a phase other than a carbide occurred in the 20 per
cent molybdenum alloys of series I when the iron content was increased
fram 4 to 10 per cent. Evidence was therefore needed to establish
whether iron was having this same influence upon the 15 per cent
molybdenum alloys.

As indicated in Chapter IV, the decarburization studies were
conducted on selected series II caompositions as well as on several
alloys of series I to observe the disappearance of carbide phases,
thereby affording a method of phase identification. The decarburi-
zation treatment consisted of heat-tfeating the alloys in a hydrogen
“atmosphere for 100 hours at 2200°F. An additional aging treatment for
100 hours at 1500°F was then given the alloys in order that a comparison
could be made between these microstructures and those obtalned after
the conventional aging treatment at 1500°F.

Camparative carbon analyses of the alloys before and after
decarburization are shown in Table X, It is apparent that this treat-
ment effectively reduced the carbon content to 0.010 — 0.01l6 per cent.

Typical microstructures are illustrated in Figures 48 through 52. All

 

 
 

TABLE X

CARBON ANALYSES OF ALLOYS BEFORE AND
AFTER DECARBURIZATION TREATMENT

 

 

 

Analyzed Carbon

 

 

Alloy : Nominal Composition (Wt %) Content (Wt %)
No. Ni Mo Cr Fe Al Mn Before After
Series I V=43 Bal 10 5 b 0.5 0.5 0.076 0.012
V-S54 Bal 10 10 10 0.5 0.5 0.068 0.016
VT-45 Bal 15 5 L 0.5 0.5 0.073 0.011
VT-56 Bal 15 10 10 0.5 0.5 0.077 0.014
VT-57 Bal 20 5 L 0.5 0.5 0.070 0.012
VT-58 Bal 20 5 10 0.5 0.5 0.069 0.012
Series II vT-87 Bal 15 T - - - 0.025 0.012
VT-92 Bal 15 7 7 - - 0.021 0.011
VT-93 Bal 15 7 10 - - 0.024 0.010

 

- 1—(9..
85

UNCL ASSIFIED
. Y-25383

Figure 48. Alloy VI-43, 10 Mo - 5 Cr - 4 Fe - 0.5 Al - 0.5 Mn -
0.012 C - Balance Ni. Decarburized in a hydrogen atmosphere at 2200°F,
aged for 100 hr. at 1500°F. Etchant: Glycera Regia. 500X.

UNCL ASSIFIED
Y-24943

Figure 49. Alloy VI-56, 15 Mo - 10 Cr - 10 Fe - 0.5 Al -
0.5 Mn - 0.014 C - Balance Ni. Decarburized in a hydrogen atmosphere
at 2200°F, aged for 100 hr. at 1500°F. Etchant: Glycera Regia. S500X.

 
UNCL ASSIFIED
¥-24933

Figure 50. Alloy VI-57, 20 Mo = 5 Cr = 4 Fe - 0.5 AL - 0.5 Mn -
0.012 C - Balance Ni. Decarburized in a hydrogen atmosphere at 2200°F,
aged for 100 hr. at 1500°F. Etchant: Glycers Regia. 500X.

. UNCL ASSIFIED
. Y.25380
:! )
/ f
¥
~ ff §
2 )
\ =
\h 7 F"’fl} t
x ~ i
W s O
: \
) 5
) ¢
i L
/- N
: P g
Figure 51.

Alloy VI-58, 20 Mo - 5 Cr - 10 Fe - 0.5 AL -
0.5 Mn -~ 0.012 C - Balance Ni.

Decarburized in a hydrogen atmosphere
at 2200°F, aged for 100 hr. at 1500°F.

Etchant: Glycera Regia. 500X.
UNCLASSIFIED
Y-25381

Figure 52. Alloy VI-93, 15 Mo - 7 Cr - 10 Fe - 0,010 C =
Balance Ni. Decarburized in a hydrogen atmosphere at 2200°F, aged
for 100 hr. at 1500°F, Etchant: Glycera Regia. 500X.

 
_88..

alloys were found to be vold of precipitated particles with the
exception of alloy VI=-58 which retained a grain-boundary phase. These
results present evidence that within the experimental conditions of this
investigation, all phases in excess of the matrix could be attributed to
carbide particles with the exception of an additional phase found in
alloys VT-58 and VT-60, both of which were 20 per cent molybdenum alloys

of series I with 10 per cent iron.

 
 

-89-

CHAPTER VI
CONCLUSIONS AND RECOMMENDATIONS

Based upon the experimental results of the present investigation,

the following conclusions and recommendations can be stated:

Conclusions

L. From the standpoint of their creep-rupture strength at 1500°F
and 10,000 psi, it was possible to conveniently group the molybdenum-
chromium-iron-nickel base alloys containing a nominal content of 0.5
per cent aluminum - 0.5 per cent manganese - 0.06 per cent carbon
according to the three concentrations of molybdenum studied: 10, 15,
and 20 per cent.

2. The principal factors affecting the strength of the alloys
within these groups were: solid-solution elements, aging reactions,
the presenée of M6C~type carbides in the microstructures, and grain
size.

3. It could be concluded from the analyses and microstructures
of these alloys that the relative strength contribution of each factor
varied between the individual groups.

4., The combined effects of solid-solution strengthening by
molybdenum and the increase in quantity of dispersed M6C-type carbides
which this element promoted in the annealed materials were the pre-
dominant factors which progressively increased the strength of the

alloys grouped by molybdenum content. The only exception noted was in

 
-90-

the case of the 20 per cent molybdenum - 7 per cent chramium - 10 per
cent iron alloy which precipitated a non-carbide phase as a conseguence
of crossing a new phase boundary. The presence of this phase in the
mlcrostructure contributed noticeably to creep-rupture strength in the
later stages of test.

5. The contribution of chromium and iron to the strength of the
alloys within the individual groups could not be established with
certainty due to simultaneous variations in other factors affecting
creep-rupture behavior.

6. Because of the scatter.in rupture life between alloys of a
group, this was not a good critérion for strength. On the other hand,
times to specified amounts of strain preceding rupture appeared more
meaningful for correlative interpretation of the data.

7. Although the studies conducted upon the "high-purity" nickel-
molybdenum-chromium ternary alloys with 15 per cent molybdenum were com-
plicated . by carbide precipitation and grain-size variations, the data
ifidicated the strengthening influence of chromium to be significant in
the range of 5 to 10 per cent, but most pronounced when 10 per cent was
present. The strengthening influence of iron was interpreted as being
insignificant when amounts up to 10 per cent were added to a 15 per cent
molybdenum - { per cent chromium - balance nickel base.

8. A general consideration of all data obtained fram this
investigation favorably supports the composition specification placed
upon the alloy INOR-8. An increase in the presently specified

molybdenum range for thilis elloy, yet keeping within the solubility

 
-91-

limits, would be made at the expense of fabricability and would result
in only slight gain in strength as shown by this work. Although the
solid-solution strength contribution of chromium is of less signifi-
cance than the contribution of molybdenum and dispersed carbilde
particles, chromium is necessary at the specified concentration to
impart oxidation resistance. 1Iron at the specified concentration is in-
significant as a solid-solution strengthener, but the introduction of
this element into the alloy through the use of ferro-alloy additions

is not objectionable.

Recommendations

1. Puture attempts to determine the solid-solution strengthening
influence of specific elements in an alloy should be preceded by a
careful consideration of methods to isolate the variable(s) of interest.
Of particular importance, as shown by this work, is the elimination of
carbon from alloys containing strong carbide formers. Purification of
the alloys by a prior decarburization treatment, similar to that
described here, would be invaluable.

2. It would be of interest to determine the chemical camposition
of the carbide phase(s) which were present in the alloys studied for
this investigation.

3. For advancing the technology of the alloy INOR-8, the
influence of heat treatment upon creep-rupture behavior should be

investigated.

 
LIST OF REFERENCES

 
..95..

LIST OF REFERENCES

Clausing, R. E., Patriarca, P., and Manly, W. D., "Aging
Characteristics of Hastelloy B," ORNL-2314 (July 1957)

Metals Handbook, American Society for Metals, Cleveland, p. 1230
(1940)

Hastelloy, Corrosion-Resistant Alloys, Haynes Stellite Company,
p. 100 (May 1957)

 

 

Douglas, D, A. and Manly, W. D., "A Iaboratory for the High-
Temperature Creep Testing of Metals and Alloys in Controlled
Environments, " ORNL-2053 (September 1956)

ASTM Standards, Part 2, Non-Ferrous Metals, American Society for
Testing Materials, p. 1443 {1955)

 

la Marche, A. E., "Pilot Plant Development of a Nickel-Molybdenum
Base High-Temperature Alloy,” First Periodic Progress Report
Under Subcontract No. 1067 Under Contract No. W-TL05 eng-26,
Blairsville, Pennsylvania, Blairsville Metals Plant of the
Westinghouse Electric Corporation, (May 1957)

Lundy, T. S. and Stansbury, E. E., "A Metallographic and X-Ray
otudy of Nickel-Base Alloys of 20 — 25 Per Cent Molybdenum
and 3 — 15 Per Cent Chramium," Report No. 2 Under Subcontract
No. 582 Under Contract No. W-T4O5 eng-26, Knoxville, Tennessee,
Department of Chemical Engineering of the University of
Tennessee (1957)

 
BIBLIOGRAPHY

 
~-99-

BIBLTOGRAPHY

Beattie, H. J., Jr. and VerSnyder, F. L., "The Influence of Molybdenum
on the Phase Relationships of a High Temperature Alloy," Trans.
Amer. Soc. Metals, 49, p. 883 (1957).

 

Monkman, F. C., Jr., Grant, N. J., and Floe, C. F., "Final Report on
Development and Testing of Nickel-Molybdenum Alloys," ORNL-1990
~ (November 1955).

Parker, E. R., "Creep of Metals," High Temperature Properties of Metals,
American Society for Metals, Cleveland, pp. 1-40 (1951).

 

Preston, 0., Grant, N. J., and Floe, C. F., "Final Report on Development
and Testing of Vacuum Melted Nickel=Molybdenum Alloys with Minor
Alloying Additions," ORNL~2181 (October 1956).

Preston, 0., Grant, N. J., and Floe, C. F., "Final Report on Development
and Testing of Air Melted Nickel-Molybdenum Alloys with Minoxr
Mloying Additions," ORNL-2520 (June 1957).

Rotherham, L. A., Creep of Metals, Institute of Physics, London (1951).

Stoffel, D. W. and Stansbury, E. E., "A Metallographic and X-Ray Study
of Nickel Alloys of 20 — 30 Per Cent Molybdenum," Report No. 1
Under Subcontract No. 582 Under Contract No. W-7LO5 eng-26,
Knoxville, Tennessee, Department of Chemical Engineering of the
University of Tennessee (1955).

 
APPENDIX

 
-103-

APPENDIX
CREEP-RUPTURE TESTING APPARATUS

Typlcal creep frames used Iin the Mechanical Properties .
Iaboratory of the ORNL Metallurgy Division for testing in gaseous
enviromments are shown in Figure 53. The main camponent of a creep
frame of this type is a leaktight test chamber which can be heated uni-
formly to the desired test temperature and in which strain and tempera-
ture measurements can be made.on a stressed test specimen. The chamber,
1tself, consists of a metal tube, water-jacketed on the ends to allow
sufficient cooling for the brass bellows and for the rubber O-rings
utilized as seals. Side ports are provided to allow for optical
measurement of the specimen extension. The heating elements are woufid
around the tube in such a way as to compensate for the heat losses at
the ends and at the side ports. External shunts are provided to allow
for adjustment of vertical. temperature gradients. Pull rods connected
to and extending through the two bellows make it possible to introduce
a load on the specimen inside the chamber. With the bottam pull rod
anchored in place, the specimen can be stressed by application of a
load to the top bellows through a lever am gnd a welght-pan system.

Four thermocouples are wired to the specimen, two at either end
of the specimen gage length, and the leads are brought out of the test
chamber through rubber stoppers to a fiunctiofizbox. One thermocouple is
used for temperature control with the use of a Leeds and Northrup

duration-adjust-type controller (DAT) and a Speedomax recorder.

 
e

U NCL ASSIFIED

 

 

Creep frames for testing in gaseous environments.

Figure 53.

 
-105-

Temperature readings from each of the four thermocofiples‘are made with
a multi-point precision indicator.

The extension of specimens during creep is measfired,optically by
means of a pair of dovetailled platinum alloy strips which are spot-
welded over the gage length. The strips are referenced so that measure-
ments can be made with a microscope through the test chamber side ports.
This optical method of measuring extensions is precise to approximately
+ 0.0003 inch. Figure 54 shows a specimen connected to the top bellows,
with the platinum extensometer and thermocouples in place ready for
assembly in the test chamber. Figure 55 1s a view of the scribed plati-
num extensometer from which the extension measurements are made. A
dial gage attached to the pull rod outside the test chamber acts as a

rough check on the micrometer microscope readings.

 
Figure 5h.
and pull rods with

Assembly of creep-rupture
the top bellows flange.

specimen, extensometer,

 

 

thermocouples,

90t

 
 

107

JNCL ASSIFIED
Y. 15839
T T .

N

 

Figure 55. Gage length of a creep-rupture specimen with

 
ORNL-252L4
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