NFSD 89-0 (June 30, 1989)
FORMS
53-55
4. _ me Smum
Report Documentation Page
2. C_mm_umt A_m_mn tin.
Effect of Boron on Intergranular Hot
Cracking in Ni-Cr-Fe Superalloys ContainingNiobium
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N/A
R. G. Thompson
University of Alabama at Birmingham
Department of Hmterials Engineering
UAB Station_ Birmingham_ AL 3529412.$mm_ *4Wecv t_.m _
National Aeronautics & Space AdministratorsWashington, D.C. 20546-001NASA/M_rshall Space Flight CenterMarshall Space Flight Center, AL 35812a sew_,ma,v,.,
June 22, 1990I. _mi,,t Oq_um_ f,4m,
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Final
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Solidification mechanisms had a dominant influence on micro-
fissuring behavior of the test group. Carbon modified the Laves
formation significantly and showed that one approach to alloydesign would be balancing carbide formers against Laves formers.Boron's strong effect on microfissuring can be traced to its
potency as a Laves former. Boronls segregation to grain boundariesplayed at best a secondary role in microfissuring.
Hicrofissuring
Liquation crackingBoron
Carbon
II. Oitmbullo_ SItltmtnl
DTASulfur
SegregationSolidification
II. _ C_a_. Iof Itm ml_t) / 20. f_cu_.v C_. Iof _ D_ilel
J(NA_A-CR-I_714) t_FF_CT _F JOkON L}N
INT_GRANULA_ HOT CRACKING IN Ni-Cr-F_SUP?_,LLOYS C._i'_TAININ_ _]IqBTU_ Fina| _eport
(Alabama Univ.) 115 p CSCL llF
NASA FAR SUPPLEMENT
N90-2611 )
Uncl _s
18-53.303-1626
https://ntrs.nasa.gov/search.jsp?R=19900016799 2020-03-19T07:56:47+00:00Z
Table of Contents
Executive Summary ................................................ I
Introduction ..................................................... 2
Procedures and Results ........................................... 2
Specimen Selection ............................................ 2
Specimen Chemistry ............................................ 3
Thermal History of Alloys ..................................... 3
Microstructural Analysis ...................................... 3
Differential Thermal Analysis ................................. 4
Gleeble Thermal Analysis ...................................... 4
Scanning Auger Microscopy ..................................... 5
Discussion ....................................................... 5
Alloy Melting and Solidification .............................. 5
Simulated Welding Thermal Cycles .............................. 7
Interfacial Segregation Effects ............................... 7
Microfissuring Susceptibility Mechanisms ...................... 8
Conclusions ...................................................... ii
Figure I - Gleeble Thermal Cycle ................................. 12
Figure 2 - DTA Runs of Commercial Alloy 718 ...................... 13
Figure 3 Solidification of Boron Alloys ........................ 14
Figure 4 Typical Range of Sulfur Segregation ................... 15
Table I - Alloy Chemistry ........................................ 16
Table 2 - Thermal History of Alloys .............................. 17
Table 3 Summary of DTA Results ................................. 18
Table 4 - Results of Scanning Auger Microscopy Peak Height
Ratios to Ni ........................................... 19
Table 5 General Electric Report of Spot Varestraint
Cracking Data on Alloys of this Study .................. 21
Appendix A - Metallography ....................................... A
Appendix A2 - SEM, Energy Dispersion X-ray Analysis
of Precipitates in Samples ......................... A2.1
Appendix A3 - TEM, Energy Dispersion X-ray Analysisof Particles in Thin-Films ......................... A3
Appendix B Differential Thermal Analysis (DTA) of TestSpecimensof Micrographs of DTASolidificationMicrostructuures .................................... BI
Appendix C Microstructures from Gleeble Weld Simulation ........ C
Appendix D Scanning Auger Microscope (SAM) ..................... D
NASA Form 1626 ................................................... 22
ii
EXECUTIVE SUMMARY
The role of boron in intergranular liquation cracking (microfissuring)is now better understood. Boron both increases the solidification
range of 718-type alloy compositions and promotes Laves phase
formation. This combination is detrimental to welding because it
increases the probability that liquid will still exist on the grain
boundaries of the heat affected zone (HAZ) when residual stress begins
to develop. This is the situation that is most likely to result in
microfissuring.
A new discovery of this study is that carbon can mitigate the
deleterious effect of boron by changing the solidification to primary
carbide formation instead of Laves. This has the effect of reducing
the solidification range and reducing the possibility of the
simultaneous existence of grain boundary liquid and residual stress.
This was consistent in the differential thermal analysis, Gleeble
simulation, and varestraint test results.
Data on the intergranular chemistry of 718-type alloys shows that boron
is not strongly segregated intergranularly. However, these results
establish that sulfur and phosphorus do segregate to the grain
boundaries. It also revealed that sulfur strongly segregates to the
carbide and Laves phase interfaces. This explains the ability of
sulfur to increase the solidification range of carbide formation.
It can be concluded that boron should be avoided if it is not necessary
for high temperature mechanical properties. Sulfur and phosphorus
should be minimized. Carbon can be increased in certain chemical
compositions to good benefit. However, excess carbides can increase
the solidification time thus they should generally be avoided.
Finally, there should be a generally applicable alloying scheme thatbalances the carbide formers against the Laves formers to eliminate
Laves formation and thus produce a microfissuring resistant alloy.
.
FINAL REPORT
"EFFECTS OF BORON ON INTERGRANULAR HOT
CRACKING IN NI-CR-FE SUPERALLOYS CONTAINING NB"
Introduction
The primary mission of this research was to define the mechanistic role
of boron in microfissuring of 718-type superalloys. Because of the
experimental procedure selected, it was also necessary to evaluate the
roles of sulfur and carbon on microfissuring.
Nickel base alloys, particularly those like alloy 718 are susceptible
to HAZ microfissuring during welding. This is unfortunate since the
718 base composition is expected for use in current and future advanced
engine designs. The need to eliminate microfissuring stems from the
facts that; it is difficult to inspect for, fatigue strength is
compromised due to the loss of crack initiation time, and it can act as
a point of stress concentration and site for hydrogen embrittlement
growth.
Previous studies(l-3) have not identified the mechanism by which
various alloy and impurity elements degrade or benefit the resistance
to microfissuring. It was proposed that the only serious detrimental
element was boron(l). This is in contrast to other studies(2,3) that
have found elements like sulfur and carbon deleterious to
microfissuring resistance. It is difficult to experimentally isolate
an element like boron which is present in concentrations of only a few
parts-per-million.
This study took the approach of using a series of eight alloys that
formed an experiment involving C, B, and S. The approach was to
examine the effect of these elements individually and in combination.
By studying the mechanisms by which these alloys influenced
mlcrofissuring, it was possible to determine the individual and
synergistic effects that they have on microfissuring.
Procedures and Results
Specimen Selection
A 2x3 factorial experiment using carbon, sulfur and boron was attempted
from a set of alloys produced by General Electric. The alloy set was
based on modifications to an alloy 718 type base chemistry. Thedesired combinations were:
B i i i i 2 2 2 2 l-low concentration
C i I 2 2 i 2 I 2 2-high concentrationS i 2 i 2 i I 2 2
Spec# 20 14 11 10 13 12 3 5Subs# 25 22 27
Although it was hoped that #ii and #14 would allow Si to substitute for
C and P to substitute for S, it was determined after some
experimentation to abandon that substitution. Instead, alloys 22 and
25 were substituted for Ii and 14 respectively. These alloys were from
a different study but they met the general requirement of high and low
S and C with a 718 base composition. Alloy 27 was used as a control to
give overlap with alloy I0 to test the validity of mixing the two alloy
groups. It was found that these substitutions had some draw backs.
First, each of the substitutes contained typical impurities or alloy
levels different from the test group. These elements, such as silicon,
promoted Laves phase and the low carbon level still caused carbide
precipitation. Thus the substitutes behaved more like a combination of
elements similar to alloy 5. They did provide some insight as
discussed in the following sections.
Specimen Chemistry
The chemistry of each alloy in the study is given in Table I.
Thermal History of Alloys
The thermal processing of the alloys is documented in Table 2.
Microstructural Analysis
Light Microscope
Micrographs of initial microstructures used for DTA, Gleeble and Auger
are presented in Appendix AI for all specimens. Metallography on
initial materials was performed using standard metallographic
procedures. Oxalic acid was used for all etchings. The micrographs
show relatively clean, equiaxed microstructure after a II00C thermalanneal.
SEMMicrographs
Micrographs at approximately 3000x are presented in Appendix A2.1 for
specimens 5, i0, 12, 13, and 14. These micrographs show typical
precipitate structures and the locations of analysis points used for
chemical identification.
SEM Energy Dispersive X-ray Analysis
EDS analysis of precipitates found in specimens I0, 12, 14, 20, 5, and
13 are presented in Appendix A2.2. The spectra are divided into four
general types:
-spectra that resemble matrix (A2.2.1)
-spectra that resemble Nb-carbide (A2.2.2)
-spectra that resemble Laves phase (A2.2.3)
-spectra that resemble delta (Ni3Nb) (A2.2.4)
Spectra that resemble matrix are characterized by small Nb peak, small
Ti peak, and major Ni, Cr, and Fe peaks in descending intensities.
Spectra that resemble Nb-carbide are characterized by a primary Nb peak
3
and a small Ti, Ni, Cr, and Fe peaks. The spectra of Laves phase have
roughly equal Nb, Ni, Fe, and Cr peaks. The spectra of delta phase
have a Ni peak larger than found in the matrix and roughly equal Nb,
Fe, and Cr peaks.
TEM Energy Dispersive Analysis
TEM-EDS analysis of specimens 12(A3.1), 13(A3.2), and 20(A3.3) are
presented in Appendix A3. A TEM micrograph of specimen 12 is presented
in A3.1. Specimen 20 was practically devoid of precipitates. The
other extreme was specimen 12 which had a multitude of precipitates
from which one area was examined in detail. From this area, 16 spectra
were taken from various particles. These spectra show several things
of importance. First, the particles are embedded in the matrix and the
spectra of all particles has a strong matrix signature (A3.1.1) Two
particles are typical in this matrix; a Nb-carbide particle identified
by the large Nb and Ti peaks (A3.1.2) and a carbide containing Nb, Hf
and Zr (A3.1.3). The Hf and Zr appear at the expense of the Ti and Nb.
Spectra in A3.1.4 is of a Nb-carbide that has little matrix embedded
with it.
Particles found in specimen 13 were very small and from the spectra in
A3.2.1 appear to be oxide inclusions.
pifferential Thermal Analysis
DTA results for all specimens are presented in Appendix B. The melting
and solidification curves contain characteristic shapes which have been
marked matrix, NbC and/or Laves. When carbon is present, the DTA
results show a characteristic shoulder that appears with the initial
stage of bulk melting and the final stage of primary gamma
solidification. This shoulder is the proeutectic gamma+NbC reaction.
A lower temperature reaction which appears on cooling is the eutectic
reaction believed to involve gamma+carbide+Laves.
The solidus and liquidus temperatures for each reaction are found by
taking the point where the tangent to the baseline meets the tangent to
the reaction curve of interest. These temperatures are presented in
Table 3.
The reproducibility of these results is presented in Appendix B2. Note
that these results are very reproducible.
Micrographs of the alloy microstructure following the DTA thermal cycle
are presented in Appendix B3. Also presented in B3 are chemical
spectra from SEM-EDS analysis. These are presented to correlate the
DTA reaction peaks with the phase responsible for the reaction.
Gleeble Thermal Analysis
Gleeble experiments were conducted at four peak temperatures (I190C,
2010C, 1230C, and 1250C) which encompass all of the liquation
4
temperatures found in these alloys. These Gleeble thermal cycles are
shown in Figure I. Samples were heated to the peak temperature and
quenched in water to nfreeze in" the microstructure for metallographic
analysis. The Gleeble results are summarized in Table 3 and
micrographs of the samples from the Gleeble tests are presented in
Appendix C.
Scanning Auger Microscopy
SAM was performed on all samples in the test group. About 400 spectra
were taken of surface chemistries of grain boundaries, cleavage facets,
ductile ruptures, carbides, Laves and other second phase particles.
Samples were prepared by cutting immxlmmxl0mm specimens from the
alloys. The specimens were then put into a hydrogen charging cell to
embrittle the grain boundaries for fracture in the SAM. The parameters
for hydrogen embrittlement are given in Appendix DI. Samples for SAM
were placed in the SAM and fractured using a slow strain rate. The
fractured surfaces were then analyzed using standard SAM techniques.
Typical spectra from each sample are presented in D2. The results from
the SAM study are presented in Table 4. Since over 400 spectra were
taken, the presentations represent typical results obtained.
In order to determine if the observed Auger chemical spectra is a
surface film or bulk chemistry, depth profiles were made by sputtering
the surface under analysis with argon ions. It can be assumed that the
spectra presented are for unsputtered surfaces unless otherwise
designated. The spot probe used to obtain chemical spectra was i to 3
microns in diameter. The largest diameter was usually used unless
particle size or other features required a smaller probe size.
Discussion
Alloy Melting and Solidification
The DTA was used to evaluate the effects of B, C, and S on alloy
melting and solidification. The method for doing this can be shown
with DTA data from a commercial alloy 718 as presented in Figure 2.
The initial microstructure is free of Laves phase but contains some
NbC. When this microstructure is heated (Figure 2A) a smooth baseline
is established until the NbC-solidus (1245C) is reached. The sample
absorbs heat during NbC melting and then the rate of this reaction
changes as bulk(gamma) melting begins. During solidification (Figure
2B), the gamma solidification reaction is encountered first followed by
a second exothermic peak associated with NbC proeutectic
solidification. The final solidification reaction is the system
eutectic which has been described as a gamma+Laves+carbide. The
microstructure verifies this solidification sequence. Note that this
718 mlcrostructure contains a significant amount eutectic which, when
heated, produces an endothermic reaction associated with eutectic
melting (Figure 2C). The NbC and gamma melting reactions are also
identified. This exercise points out that alloy 718 melting reactions,
that is the heating of the initial microstructure, are dependent on the
initial microstructure. This is clear from Figure 2 where the Laves
peak is present during the heating segment of the DTArun only if Laveswas present in the initial microstructure. However, the solidificationreactions which occur during cooling are dependent on chemistry and
solidification rate as seen in Figure 3. The DTA is capable of
evaluating this behavior.
Interstitial Free (#20)
The interstitial free alloy had the smallest freezing range, 48C, due
to the lack of carbide or eutectic solidification reactions (BI.6).
Single Level Effects B(13) S(25) C(22)
The effect of a single interstitial element can be dramatic, as with B
(alloy 13) which acts as a strong Laves former and increases the
solidification range to 174C (a factor of 3.5 increase over alloy 20)
(BI.5). It should be noted that phosphorus (alloy 14) has a similar
effect (BI.9). Carbon (alloy 22) also has a dramatic effect in that it
produces carbide precipitation which increases the solidification
range, gamma liquidus to carbide solidus, to I16C (note that this
solidification range did not include the Laves reaction which
presumably appeared because the alloy contains other interstitials)
(BI.7). Sulfur appears to have a dual effect by accentuating both the
carbide and Laves reactions (alloy 25 had residual carbon and Laves
formers) (BI.8).
The alloy set had a varying Nb content that may have had an effect on
the magnitude of the Laves reactions. For example, alloy 20 had 4.4wt%
Nb while alloy 13 had 5.4wt% Nb. While B appears to be a strong Laves
former the observed behavior is probably influenced by the increase in
Nb concentration. However, in defense of boron's influence on Laves
formation it can be noted that boron not only increased the amount of
Laves formation, it also changed its morphology from that of a
gamma-Laves mixture(which is normally reported in alloy 718) to a
blocky, unmixed Laves precipitate.
Binary Interactions BC(12) SC(I0,27) SB(3)
Binary interactions proved very important. The BC interaction was most
dramatic with carbon causing a complete solidification change from
strongly Laves eutectic in the boron alloy (13) to completely carbide
(BI.4) in the BC alloy (12). This observation should be tempered with
the fact that BC alloy had the low Nb concentration(4.4wt%) which could
contribute to the loss of Laves formation. The solidification range
for BC was very small, only 62C. The SC interaction produced strong
carbide formation and a microstructure like the BC alloy. The
solidification range, 60C, was the same as that of the BC alloy (BI.3).
(Alloy 27 which contained other interstitials had the same
solidification range as #22 when considering the gamma liquidus to
carbide solidus range). Alloy 3, used as the SB binary interaction
alloy, actually contained phosphorus also which, based on P behavior,
probably accentuated the Laves eutectic behavior (B and P gave
identical DTAresults). However, the solidification range for alloy 3was smaller than for either P or B alone.
Ternary Interaction BCS(5)
Alloy 5 actually contained both P and Si along with B, C, and S. Thesolidification range was large as would be expected for Lavessolidification. The ternary interaction was similar to that found for718 base alloys with typical interstitial levels.
These results show that the interstitial chemistry of alloy 718 can bethought of in terms of carbide or Laves formers. The Laves formerstend to extend the solidification range while the carbide formers havelittle effect on it. However, the carbide formers can suppress theLaves formation as was the case with the BC alloy and therebydramatically reduce the solidification range. Commercial 718 alloybehaves like alloy #5 in this study in that the commercial chemistrycannot offset the multitude of Laves formers and a duplex carbide/Lavessolidification is obtained with the attendant large solidificationrange (Figure 3).
Simulated Welding Thermal Cycles
Gleeble simulation of the weld heating cycle was used to examine the
relationship of the DTA melting reactions done at a slow heating rate
to that of the weld HAZ which is a relatively rapid heating rate. TheGleeble results do indeed reinforce what was discovered in the DTA
work:
-B, P, and SB(P) alloys(13,14,3) initiate liquation below
I190C. These were all strong Laves forming alloys.
-BCS alloy(5) initiated liquation at about 1210C.
-CS and BC alloys(lO,12) both initiate liquation between
1210C and 1230C. These are the carbide forming alloys.
-Alloy 20 initiates liquation at 1250C.
The Laves formers, in the absence of carbon, produce a wide melting
range and presumably an even wider solidification range. The addition
of carbon apparently raises the eutectic thus giving a smaller melting
and solidification range. If enough carbon (relative to Laves formers)
is present, then the eutectic is completely suppressed and melting and
solidification ranges are similar to the interstitial free alloy.
Interfacial Segregation Effects
The distribution of the interstitial elements could have significant
effects on the solidification and melting behavior of the alloys. If
the interstitial elements are strongly segregated on the boundaries or
at the phases which melt, they can more strongly influence the behavior
than if they are dilutely and homogeneously distributed. Also, whether
or not the interstitials partition to the solid or liquid during
solidification strongly affects the solidification of the alloy.
7
The general finding of this study was that when the alloy was doped
with an interstitial, that interstitial appeared as a grain boundary
segregant. Thus B, S, and P all appeared on the grain boundaries when
they were present in the alloy at high concentration. However, sulfur
was always found at the NbC-matrix interface in huge amounts(lO to 40
times greater at the precipitate interface than on the grain boundary)
as seen in Figure 4. Sputtering to remove the surface layer removed
the sulfur concentration thus proving that it was not part of the bulk
chemistry of either the matrix or carbide. Phosphorus was also found
to be more often associated with grain boundary segregation than either
sulfur or boron. Boron was found on the grain boundaries but never
associated with any precipitate. However, sulfide particles were often
found on fracture surfaces.
These findings provide significant new information about the possible
role of boron on mlcrofissurlng. This can be evaluated by comparing
the BC alloy (12) and the B alloy (13). Note that both alloys are
doped with B and both have about the same B level on the grain
boundaries. As will be noted later, the BC alloy exhibited little
cracking (using the total crack length as the measure of cracking)
compared to alloy B. Since there was a large difference in the
cracking behavior but little difference in the grain boundary boron
concentration, intergranular boron must not be the cause of cracking
susceptibility. This strongly suggests that the solidification range
has a much greater effect on microfissuring than does the intergranular
segregation level exhibited by boron.
The sulfur concentration at the carbide-matrix interface has the most
potential to dramatically affect liquation and thus microfissuring.
The reason for this is the large concentrations of sulfur possible at
the precipitate-matrix interface. Furthermore, it appeared that sulfur
was also heavily segregated at the Laves-matrix interface although this
was difficult to confirm due to the eutectic nature of the Laves phase
and its general Auger spectrum that looks like matrix. The obvious
implication is that S partitions to the liquid phase during
solidification and is eventually rejected during the final
solidification stages, which are either carbide formation or eutectic
formation. Due to the common association of S with the carbide
interface, it must be concluded that some mutual chemical attraction is
present. The potential effect of this S segregation could be on
controlling the shape of the carbide or on the liquation behavior of
the carbide-matrix interface. However, the alloys containing sulfur
were not found to be exceptionally prone to microfissuring nor did they
exhibit eutectic behavior in the Gleeble rapid heating tests.
Microfissuring Susceptibility Mechanisms
Table 5 presents microfissuring susceptibility results from spot
varestraint tests on the alloys of this study. Although both the total
crack length and average crack length are reported, all correlations
and references to cracking in this report refer to the total crack
length measurement. Past experience of this researcher and others has
8
been that total crack length correlates best with cracking
susceptibility. In the General Electric test group, the strong Laves
formers in the absence of carbon gave the greatest level of
microfissuring. Just as Laves appears to strongly promote
microfissuring, carbon appears to improve microfissuring resistance.
It should once again be noted that alloys containing carbon also had
the lower niobium content which could also contribute to reduced Laves
formation. Sulfur gave no strong indication as to its effect on
microfissuring from the cracking data.
Effect of Boron on Microfissuring Susceptibility
Boron promotes a long solidification range through promotion of Laves
formation. The boron containing alloy(13) had the lowest melting point
in the Gleeble weld simulation tests which indicates it ability to
promote intergranular liquation during rapid heating. Alloy 13 also
had the most microfissuring susceptibility. However, such behavior is
not necessarily unique to boron. It was also observed in the behavior
of phosphorus additions(alloy 14). Although researchers at General
Electric found that only boron significantly promoted microfissuring
based on statistical results of many compositions, this research shows
on a specimen by specimen basis that the boron effect on microfissuring
is not unique. Any combination of elements which cause a long
solidification range and promote a significant amount of low-melting
eutectic phase will have similar microfissuring susceptibilities.
Boron's contribution to microfissuring susceptibility can be negated
through elimination of the low melting phase which it promotes. Such
was the case with alloy 12 where the carbon level was increased and
the Nb level reduced so that the Laves solidification was suppressed.
The result was that microfissuring was also suppressed even in the
presence of a large boron concentration.
Effect of Solidification on Microfissuring Susceptibility
The solidification temperature range appears to be a good first order
approximation for microfissuring susceptibility. This finding should
be modified to also encompass the amount of eutectic formed at the
low-temperature end of solidification. Examples of short
solidification range and low microfissuring are alloys 20, 12, and I0.
Alloy 5 is an exception, but it contained O.IC which modified its
solidification such that it had a very small Laves DTA peak. Although
technically it had a large solidification range, the amount of Laves
formed may have been insignificant in terms of microfissuring. The
large carbon levels (possibly combined with smaller Nb contents) used
in this study appear to alter the solidification path sufficiently to
reduce microfissuring, probably by reducing the volume fraction of
Laves phase. This was most noteable when the Laves was completely
premoved from the solidified microstructure.
A large solidification range with significant Laves present, as
indicated by the size of the DTA peak, produced the most
microfissurlng(total crack length). This is evidenced in alloys 13, 14and 3.
Effect of Segregation on Microfissuring Susceptibility
A dramatic result of this study was found in alloys 12(CB) and 13(B).
Both had about the same level of grain boundary B segregation but alloy
12 had very little mlcrofissuring(total crack length) similar to the
clean alloy (20). This must be a reflection of the dominating effect
that solidification has on microflssuring as compared to the observed B
segregation to grain boundaries.
Sulfur appears to have little to do with the microfissuring levels of
this study. That might be due to the fact that most of the sulfur was
either tied-up as a sulfide or segregated to the carbide/matrix and
eutectic/matrix interfaces. This may be why sulfur appeared to have no
effect. Sulfur may really just have the same effect on all specimens
because it was found in both samples with and without sulfur doping.
Interstitial Concentrations
The absolute concentrations of interstitials used in this study are
well above those normally encountered in commercial alloy 718. This
may be the cause of certain differences in trends between the General
Electric test group and the auxiliary test group (22, 25, 27),
especially with respect to the carbon concentration. The auxiliary
test group all showed a tendency to Laves formation which is probably a
result of elements like Si which increase Laves formation. The large
carbon concentration in conjunction with the lower Nb content of the
carbon-bearing alloys in the General Electric group probably changedthe solidification relative to the commercial carbon level. This
allowed the GE alloys to solidify completely as primary carbide and
completely avoid Laves formation.
10
Conclusions
.
.
Boron increases microflssurlng by its potency as a Laves former
and the resultant long solidification range that Laves-forming
alloys have.
Boron's behavior is not different in this respect from an element
llke P which gave strikingly similar results to B in both the DTA
and microfissuring tests.
. The segregation of boron to grain boundary surfaces is evident in
alloys doped with 0.01wt% B. However, this level of segregation isnot sufficient to offset solidification effects associated with
primary carbide solidification compared to solidification to a
Laves eutectic.
4. Boron is not unique in its ability to promote microfissuring or in
the mechanism through which it promotes microfissuring.
.
.
Carbon, in large concentrations (>0.1%), can significantly alter
the solidification behavior of the 718-type composition and
completely reverse the effect of a Laves former like boron. This
was observed in a boron alloy with 5.4wt% Nb and a boron-carbon
alloy with 4.4wt% Nb.
Sulfur segregates strongly to carbide and eutectic interfaces
whether or not it is present as an intentionally added dopant.
This probably explains why sulfur additions did not strongly alter
microfissuring data.
. Results from Gleeble simulation of the welding heating cycle
provided information to verify that results from DTA studies were
applicable to microfissuring studies.
. Alloy design for minimizing microfissuring should consider various
options such as:
-extra low interstitial compositions,
-high carbon compositions,
-compositions which balance carbide formers against Laves
formers, and
-compositions which can be made free of second phase
precipitates which liquate and produce a liquid with a large
solidification range.
11
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....... _ ................. _ ...... J ....... _ .... _ .... 4 ...... 7++ ..... _+--_ .................... _ --I ..............J---__+_--...... _.:--_--=I...............-i..........: : |_:_._ "++....._..................., ....
P-- ..... -t .... _ ...... _ ....... 4-- .................. C" -.:_ -: .................... ] ..... -' ........._-- ..... _ -----_-- -4--- -_........................ -_--:...... . • i ..... ] ............ +,--q.... ........!....... . ,--,--+---+2++0¢+-::-, +-::-,
+- --'::- ____ __ ____...t ____- -::-L :'_-::: .... 'i..................... :I _ ......... - ...... '_- _--.............-- ' I --"------:I- _ "_:-_---:::I .................... +.....................:.-_--- -:T+:._._--___--_ -_--- ------'-.-_:.--+'_- _ :'-_--_L::-- -:__ -----:- T--_" I ............... + .... 4--- +-+ .............. ] .- :. : . • : _
.......................................................... ,....._._.+OC ......
• + ..... -. ................................. ; ...... 1
+ ' I ........ , ...... 7 ....... 1.......... 7--T--1 ...... t.......... //+/_--+---"i. .... -i......... " - : ++' ' :---:_:--- -: .... ::.+1............ _...................... I ....... + ....... :-- -_-------- --_
' ........ :-I ................. I ....... i.... -1 ........ _ ....... t ///i - " I ..... + : : Ip ........................ +i.... _............ +................ +' .......---__----:f...... : +....... + .............. i ............. //// ++...................... J. Z_U:Z--I--=Z-:--+_-_::----=Z----:+ ::-::---_:ZZ:Z._='--IZZT-]+_-=--.]II-_-: //]/. " + _:_: .L: I : . I : + + Z + 7 : ] . . ; " +
L___ ++t ......... -----+ _I .... _ .... u ............... u ...._--- --- _ _ ! ---- --_++__ ................................. J ..... - ................................ 1
t i ...... +......... i........... t.... I ......... F// -:-I: " • ...... i .... i ...... --::--: --q-_-__-:_--- _--:-----__::-:-_:I.:__::-:_:;_:<::- J .... /// + . ! ......... t ......... --I
,-- ........ , ............................... +......++ +,,,, + • , +_-------:I:..._: -- -- --'--:- +.... ::__+1 ..... :. _:_::-._ _: _ VZ/ .__ ', ...... + -
, " -7 ..................... *', - ! .......... .._--.............. + ........................ I.:.-:::-___i.-:+ : 4+E ' . .:__ : .- , • .+ ........-- ::-it-+.....:::-::-:::::--+_-:-:-::-J+:::-:_--=::=::--ilf*----+::::+ _ii-:-i++:! " - :--: ++'" " i : ]_:---i ..... I ....... -i.......... +......... ,/' .... I. . + . i i.-: " •_:--__---:.:_-iT_----3__---Y-_--_:_-3-1_::i:-:_:-.:::---_iV+++-:--_-__-]-+i::-I+::.+.++ :++-:--+-:+ + +: ; "
..... ++.... I .... 4+ .............. :J_i'- -4 _, " ...... I ....... '....................... j ............ -- --+ ......... H
_. , -,--i + ........... , ...... • • , i _i+ I ' i I
- i, ................ + ...... / - _:l-c:- - :i ..... I ....... I" - - + ]
--__-_ ....... _--- ...... '1- --: _.:z +J, :-:----; -- :: + - : - - • _ .... + : t
!........ • ...... ; --- I .l!_!. l---::.; -_----_ " :'-',,' _ _
•., O_u i _. ! ] + .: i " ' . - •
--_ -t- .... -_ -:: ...... _./ -:-\_- : " - + l_ +..... :.-- 7- . :.:: ,'.---- : : ! -: -: i 7:7 ! _i..... :- " - _ i i " :-:-- i .... '
• - + - / ' i/-.:__-.__L. .: - -- :_ .-: ':-:--+:-- -t _ I.:.'. - . -; .... - 7.- 7 : +-:. i 7--7---7 T . -- ": .... -:- .... ! -- 7-:'i -/ _ - 7 " " +t ; - " ' ',
' - ..... _+--+---:"/+ - -i ' : ' ' .............- ..... :. :- . .::. -_ _-v:-t " . .. , .... I
:+_:_i- _, ---: - : :--:-:Q------:-:-----I _-: -;: ......:_t__ -- " ' ;...... , --i7 +........ ! ..... +. .... 1 ' : ' +.... - " : it I - ._: ---1 -- " -:----: " t : ; " . . . .........
'- :--- / ;- -i --- : :-- i : : ' " '-i - i'i ....... i .......... ' !
" -:-;_.-.- : li.i:v-:-:) !::: _ -:---:: :-i : " _.... i " ! " '__: _ --- -#. ---.u__ f____._ ....... _ : -_ • . ,
i " ;
....... :- : . + ....... -- r --"
i -" ! : +'+- + : ......_+_- :-+ ---+. +--T-TM E_u.(_6ee.)_
Figure i. Thermal cycle used in Gleeble machine for simulation of HAZheating.
............... . i -::! :i--iii-::i12
i , i ..: -:! . . i i
,H,]
_iii;_Iiii
_iiii
2:i
" !
_J
ilIl
DTA
1'1' _1
J! ,,1:1,iil_i|l
',, !il:tl_q:'ll
]il II
lil...l
:iii
HI& I
; !-4_
"'[
: I
,,.!
!il,!" I
DT,r
Figure 2. DTA runs of commercial alloy 718 melted with only carbide
(top), solidified (middle), and melted with Laves and
carbide which resulted from solidification in DTA (bottom).
Accompanying mlcrographs of mlcrostructure prior to first
melting (top) and after solidification in DTA (bottom).
13
ORIGINAL PAGE IS
OF POOR QUALITY
it.t
on cooling
matrix
tn t.tt
Figure 3.
on coollng,=
7 matrix
maCrix_
on cooling
ORIGINAL PAGE IS
OF POOR QUALITY
ComposiCe of solidification of Boron alloys 13 (top),
commercial alloy 718 (middle), and 12 (bottom). Note how
carbon eliminates the Laves phase.
14
4J_4
@&J
0
i "
-50O
tO0
HO
0
-2OO
EO0 I_00
eV-_
Figure 4. Composite of alloy 27 showing typical range of sulfur
segregation. Far left is "clean" boundary after sputtering,
middle is typical grain boundary sulfur segregation in
doped alloy. Far right is huge sulfur segregation at a
carbide�matrix interface typical of all samples.
15
TABLE 1
ALLOY CHEMISTRY
Alloy
3 (BS(P))
5 (13Sc +)
1o(cs)
27(cs+)
12(cl_)
13 (n)
14 (P)
20 (none)
22 (C +)
25 (S +)
Ni
13al
Bal
Bal
53.18
Bal
Bal
Bal
Bal
53.18
53.77
C
LAP
0.1
0.1
0.065
0.1
LAP
LAP
LAP
0.065
0.023
Si
LAP
0.35
LAP
<0.01
LAP
LAP
LAP
LAP
<0.01
<0.01
S
0.015
0.015
0.015
0.09
LAP
LAP
LAP
LAP
0.003
0.09
P O_(Nb) B
0,015 5.4 0.01
0.015 4.4 0.01
LAP 4.4 LAP
<0.001 5.07 0.003
LAP 4.4 0,01
LAP 5.4 0.01
0,015 5.4 LAP
LAP 4.4 LAP
< 0.001 5.07 0.003
<0.001 4,39 0002
Fe
17
17
17
Bal
22
17
22
17
Bai
Bal
No
3.5
3.5
3.5
3.15
2.5
3.5'
3.5
2.5
3.15
3.15
Hf
LAP
LAP
0.1
0.1
0.1
LAP
LAP
Zr
0.I
LAP.
LAP
<0.01
0.I
'LAp
LAP
LAP
<0.01
<0.01
Cr
20
20
20
19.17
20
20
20
20
19.17
20.46
AI
0.5
0.5
0.5
0,5
0.5
0.5
0.5
0.58
0.53
Ti
1.0
1.0
1.0
1.0
1.0
1.0
1.0
I.II
1.09
16
TABLE 2
THERMAL HISTORY OF ALLOYS
First Cold-roll Followed by Anneal
Alloy No. Starting Thickness Final Thickness Reduction (%)
13 0.2404" 0.1850"
20 0.2437" 0.1876"
12 0.2534" 0.2154"
05 0.2480" 0.1882"
14 0.2480" 0.2169"
03 0.2542" 0.1895"
i0 0.2395" 0.1801"
ii 0.2534" 0.2154"
23 0
24 7
15 0
24 0
12 5
25 0
24 8
15 0
Annealed at IO00C for 30 min. (in a box furnace) followed by air cool.
Second Cold-roll Followed by Anneal
Alloy No. Starting Thickness Final Thickness Reduction (%)
13 0.1790"
20 0.1825"
12 0.2135"
05 0.1876"
14 0.2173"
03 0.1904"
I0 0.1793"
II 0.1871"
0 1506"
0 1544"
0 1894"
0 1504"
0 2024"
0 1691"
0 1484"
0 1307"
15 9
15 4
Ii 3
19 4
6 9
ii 2
17 2
30 1
Annealed for I hr. at IO00C and air cooled for 45 min. prior to 3rd cold
roll.
Third Cold Roll of Some of the Alloys to Level Out Sample Thickness
Alloy No. Starting Thickness Final Thickness Reduction (%)
03 0.1638" 0.1609" 1.77
12 0.1887" 0.1767" 6.36
14 0.2010" 0.1853" 7.81
All the rest were not subjected to the 3rd cold roll.
Final Anneal for Recrystallization
I hour at 1000C in box furnace followed by air cool, plus
i hour at II00C in box furnace followed by air cool.
Sample Preparation for Gleeble
3 specimens were cut out of each anneal sample and vacuum sealed and then
annealed at II00C for I hour followed by water quench.
17
SOLIDIFICATION
TABLE 3
SUMMARY OF DTA RESULTS
Alloy 7L i quidus 7sot i dus _TbCst ar t N]:)Cf i ni sh Eutect ic
3 (BS (P)) 1300 - 1140
5 (BSC+) 1324 - 1284 1145
i0 (CS) 1325 - 1298 1265 -
12 (CB) 1332 - 1320 1270 -
13 (B) 1319 - 1145
14 (P) 1338 - 1150
20 (none) 1325 1277
22 (C+) 1327 - 1279 1209 1102
25 (S+) 1330 - 1235 1191 iii0
27 (CS+) 1327 - 1278 1212 1106
Commercial 718 1312 - 1243 1140
MELTING
3 (BS(P)) 1320 1195
5 (BSC+) 1324 - 1290 1250
I0 (CS) 1333 - 1300 1265
12 (CB) 1338 - 1318 1269
13 (B) 1321 1225
14 (P) 1328 - -
20 (none) 1338 1271 -
22 (C+) 1339 - 1293 1246
25 (S+) 1345 - 1282 1236
27 (CS+) 1340 - 1293 -
Commercial 718 1331 - 1270 1245
1180
1159
1160
1159
AT s"
160
179
60
62
174
188
48
225
220
221
172
125
74
68
69
96
148
67
180
185
181
86
* AT s - Solidification Temperature Range (°C)
* AT M - Melting Temperature Range (°C)
18
Table 4. Results of Scanning Auger Microscopy Peak Height Ratios to Ni.
Peak Height Ratios*
Specimen/
Area P* Ti O Fe Ni/_,8 .981 q/9 So7 C_9 _ 898 Appendix
#_/3g.b.matrix
.23 .06 .16 .03 .08 .19 .15 .07
.00 .02 .13 .00 .03 .17 .12 .07
#5g.b. .22
rg.b. X .00
matrix .00
carbide .00 I
carbide x .00
carbide + .00
00 .14 .00 .06 .15
00 .09 .00 .06 .15
00 .04 .00 .00 .12
90 1.17 .00 .08 2.29
18 .85 .00 .07 I.i0
O0 .83 .00 .ii 1.05
.16 .62 .24 1.00
.21 .48 .21 1.00
#1___og.b. #i .00
g.b. #2 .00matrix .00
carbide .31
.04 .12 .02 .05 .18 .15
.II .13 .02 .06 .24 .16
.05 .08 .00 .03 .II .09
.63 1.69 .00 .17 2.92 1.08
#1__/2g.b. #I .00
g.b. #2 .00matrix .00
.00 .20 .04 .05 .II .07
.00 .16 .03 .05 .09 .06
.00 .ii .00 .05 .I0 .Ii
#1__/3g.b. .00
matrix .00
.03 .12 .03 .05 .06 .06
.03 .12 .00 .04 .i0 .08
#2___og.b. #i .06
g.b. #I x .06
g.b. #2 .00matrix .00
precipitates .00
.04 .13 .00" .08 .24 .09
.00 .i0 .00" .04 .07 .30
.00 .I0 .00" .03 .09 .06
.00 .I0 .00" .04 .12 .09
.76 .08 .00" .02 .12 .23
#22.1
g.b. #2
g.b. #2 xmatrix
carbide
eutectic
.00 .02 .13 .00
.00 .00 .15 .00
.00 .00 .12 .00
.00 .00 3.86 .00
.00 1.01 .24 .00
09 .06 .i0
26 .07 .13
12 .06 .28
86 .64 .91
67 .39 .60
64 .35 .24
57
55
54
62
5O
5O
25 1.00
24 1.00
24 1.00
29 1.00
29 1.00
27 1.00
.05 .09 .53 .23 1.00
.06 .06 .58 .23 1.00
.08 .24 .67 .30 1.00
.82 1.77 .55 .31 1.00
.06
.06
.09
.06
.07
.15 .59 .23 1.00
.06 .59 .25 1.00
.16 .62 .29 1.00
.07
.I0
.06
.09
.15
.09 .64 .28 1.00
.03 .59 .27 1.00
.04 .07 .ii .07
.07 .00 .21 .07
.02 .09 .i0 .09
.00 5.42 1.49 .98
.07 .i0 .16 .ii
.48 .68 .27 1.00
.13 .63 .26 1.00
.19 .59 .26 1.00
.26 .64 .26 1.00
.94 .45 .24 1.00
.41 .68 .30 1.00
.05 .57 .29 1.00
.12 .62 .29 1.00
.45 .62 .36 1.00
.40 .56 .30 1.00
D2.2 1
D2.2 IB
D2.2 2
D2.2 3
D2.2 3B
D2.2 3C
D2.3.1
D2.3.2
D2,3.3
D2.3.4
D2.4.1
D2.4.2
D2.4.3
D2.6.1
D2.6.1B
D2.6.2
D2.6.3
D2.6.4
D2.7.1
D2.7.1B
D2.7.2
D2.7.3
D2.7.4
19
Specimen/Area P* S Nb B* Mo C N* Ti
#25.1
g.b. #I .00 .01 .II .00" .06 .06 .07 .06
g.b. #I x .00 .02 .12 .00" .03 .00 .28 .07matrix .00 .02 .04 .00" .03 .06 .07 .05
carbonitride .00 4.00 .68 .00- .I0 1.24 1.32 .85
eutectic .00 .14 .16 .00" .05 .13 .16 .14
O Cr Fe Ni
.23 .63 .28 1.00
.Ii .60 .31 1.00
.25 .63 .33 1.00
.88 .46 .39 1.00
.30 .57 .28 1.00
927.3
g.b. .00 .14 .20 .00 .08 .15 .14 .12 .36 .70 .33 1.00
g.b. x .00 .00 .28 .00 .08 .06 .25 .09 .05 .71 .33 1.00
matrix .00 .00 .15 .00 .05 .06 .14 .09 .20 .56 .27 1.00
carbide .00 4.61 1.88 .00 .12 3.42 2.64 2.07 1.08 .63 .40 1.00
#27.4
g.b. .00 .09 .05 .00 .02 .i0 .05 .06 .14 .59 .25 1.00
matrix .00 .00 .22 .00 .I0 .08 .14 .I0 .23 .84 .34 1.00
carbide .00 .00 1.33 .00 .i0 1.72 .57 .41 .58 .62 .32 1.00
eutectic .00 2.12 .81 .00 .09 .80 .77 .60 .41 .48 .24 1.00
carbide/
interface .00 1.81 .97 .00 .07 1.34 .80 .58 .57 .59 .29 1.00
Appendix
D2.8.1
D2.8.1B
D2.8.2
D2.8.3
D2.8.4
D2.3Acast.l
D2.3Acast.iB
D2.3Acast.2
D2.3Acast.3
D2.3B(I093).I
D2.3B(1093).2
D2.3B(1093).3
D2.3B(I093).5
D2.3B(1093).4
* element peak height relative to Ni peak height (element/Ni)
x sputtered
+ resputtered
* Because P overlaps with a small Nb peak, it must have a ratio of about i:i with
Nb to be significant.
* B overlaps with chlorine spectra
* N overlaps with Ti spectra
20
Alloy
3 (BS (P))
5 (BSC +)
10 (cs)
12 (CB)
13 (B)
14 (P)
20 (none)
TABLE 5
General Electric Report of Spot Varestraint
Cracking Data on Alloys of This Study
Total Crack Length (mm)
Homogenized at I093C
Average Crack Length (mm)
Homogenized at I093C
6 8
2 3
2 6
3 6
i0 1
9 9
13
0.97
0.36
0.21
0.89
1.21
0.51
0.33
22 (C +)
25 (s +)
27 (CS +)
From UAB Study
18.2
19.4
21.84
0.72
0.66
21
Appendix A
Netallography
A1 Netallography of Initial Microstructures
AI.I Alloys 3 (BS) and 5 (BSC)
AI.2 Alloys i0 (CS) and ii (Si)
AI.3 Alloys 12 (BC) and 13 (B)
Al.4 Alloys 14 (P) and 20 (none)
Al.5 Alloys 22 (C+) and 25 (S+)
AI.6 Alloy 27 (CS+)
A2 SEN-EDS of Typical Phases Found in the Initial Microstruc_es
A2.1 Matrix
A2.2 Nb-Carbide
A2.3 Laves Phase
A2.4 Delta Phase
A3 TEM-EDS of Typical Phases Found in the Initial Microstructures
A3.1 Alloy 12 (BC)
A3.2 Alloy 13 (B)
A3.3 Alloy 20 (none)
A
+
20Fro,.
Top - Microstructure of Alloy 3 (BS) after II00C homogenization.
Bottom - Microstructure of Alloy 5 (BCS) after IIOOC homogenization•
AI. I
Top - Mierostructure of Alloy i0 (CS) after llOOC homogenization.
Bottom - Microstructure of Alloy II (Si) after llOOC homogenization.
AI. 2
Top - Microstructure of Alloy 12 (BC) after IIOOC homogenization.
Bottom - Microstructure of Alloy 13 (B) after II00C homogenization.
AI. 3
20Fro!
Top - Mlcrostructure of Alloy 14 (P) after ll00C homogenization.
Bottom - Microstructure of Alloy 20 (none) after I100C homogenization.
AI. 4
50p.m
50p.m
Top - Microstructure of Alloy 22 (C+) as cast.
Bottom - Miicrostructure of Alloy 25 (S+) as cast.
AI. 5
Microstructure of Alloy 27 (CS+) as cast.
AI. 6
Appendix A2
SEN - Energy Dispersive X-ray Analysis of
Precipitates in Samples
- Spectra that resemble smtrix A2.1
- Spectra that resember Nb-carbide A2.2
Spectra that resemble Laves Phase A2.3
- Spectra that resemble delta (Ni3Nb) A2.4
(See SEN micrographs for location of spots)
A2.1
©/m
,--4
18 I IQ Ti_. 3" 14 II_
SEM-EDS of spot M from Specimen 12 which was typical matrix.
62.1
ORIGINAL PAGE i3OF POOR QUALITY'
,.i,,J
dl.¢..
t,l.b
",lb Cr
Fe
[4 16
iI'! i
tNi
18 i_e- -- i [2-
E,_.g_-"
SEM-EDS of spot 2 in Specimen 5 typical of NbC particles.
A2.2 ORIGINAL PAGE IS
OF POOR QUALITY'
_4
2 Ie: I 10 lIZ r].4::-_Yl6 "_- 118
SEM-EDS of spot I in Specimen 5 typical of Laves phase.
A2.3
-4J
.v -
r,ib
_,I¢'
FL
NI e
INb r:
12 I_ 16
Nb Nb
18 110 I !2 114: 1i6 i i;3
F_ner_ ---
SEM-EDS of spot 3 in Specimen 12 typical of delta phase.
A2.4
Appendix A3
TEN - Energy Dispersive X-ray Analysis ofParticles in Thin-Film
- Specimen 12.
- Specimen 13.
- Specimen 20.
TEMmicrograph and schematic EDS
spectra of particles 1-16
Schematic of microstrucCure EDS
spectra of particles 1-4
EDS Spectra of one small particle
A3.1
A3
Particle A
Photo B
A3.1 Areas used for TEM-EDS Analysis of Specimen 12.
A3.1.O ORIGINAL PAGE IS
OF POOR QUALITY
t°_
i
12 14 16 i 8 11[_._ T 12 _ 14 I16 I 18
_%_ --
TEM-EDS spectra from area typical of matrix in Photo A, Specimen 12.
A3.1.1
Hi
,t
Nb, Mo ZrNb
TEM-EDS spectra from area typfcal of partfcles 1, 3 6, 8 9, 12 and16 in Photo A, Specfmen 12. ' ' '
A3.1.2
°-
_q
't
C_- Fe.
I 2 J 4 IE,"_-T5
I E: "[1_: 112 ] 1,I ]16 [1_
TEM-EDS spectra from area typical of particles 2, 5, 9, I0, 13, 14, and
15 in Photo A, Specimen 12.
63.1.3
I
E r,_r_ ..-,-
TEM-EDS spectra from particle A of Photo B, Specimen 12.
A3.1.4
c
c_4
CF
0 Ca
12 14
Ni
Cr F_I Ni
16 I 8
Ener_
P
14b Mo
7-1o....... fi-_F-'77"_.4 "'-,,._,--'_-
TEM-EDS typical of small particles found in Specimen 13.
A3.2.1
i
coC
C
C'_"
,l/,
IZ 14
f_
i | |
I 6 I 10 112 114
Cq,_
18
Ener_ -_
116 118
TEM-EDS typical of matrix and contrast areas found in Specimen 13.
A3.2.2
D_DC_
JL !
JLJI 2 i4 i6
L../',,I ....
J8 128
Ener_ --.-
112 116 118• |
TEM-EDS typical of particles found in Specimen 20.
A3.3
Appendix B
Differential Thermal Analysis (DTA) of
Test Specimens withKtcrographs ofDTA - SolidifiedNicrostruct_res
Appendix B1BI.0
BI.I
BI.2
BI.3.1
BI.3.2
BI.4
BI.5
B1.6
B1.7
B1.8
B1.9
B1.10
DTACurves and Micrographs
DTAThermal Cycle
Alloy 3 (BS)
Alloy 5 (BSC+)
Alloy 10 (SC)
Alloy 27 (SC+)
Alloy 12 (BC)
Alloy 13 (B)
Alloy 20 (none)
Alloy 22 (C+)
Alloy 25 (S+)
Alloy 14 (P)
Alloy 718 comaercialwrought
Appendix B2 DTA Reproducibilit-y; Alloy 718 Connercial Wrought
B1
C_
BI.O
Wa
F
WW"jOwOIF.|.ll_
II
o._- .m m.w _ .m
K
1,4
H
matrix
o6o.m o_ .oo m_.oo m_ .am me_o.oo mr#o.u mlJ.oo m#,.oo
DTA of heating (melting) (see initial microstructure in Appendix AI)
and cooling (solidification) of Specimen 3 (BS).
BI. i
an |.go
o.a
matrix
f
Laves
x
NbC
.oo mio.oo m.oo om_.oo ouou.oo ooiooo mio.w nu_o.oo ,m,_,.u
DTA of heating (melting) (see initial mlcrostructure in Appendix AI)
and cooling (solidification) of Specimen 5 (BCS).
BI.2
ORIGINAL PAGE IS
OF POOR QUALITY
m |.D
wwE
wo
O.N
7 matrix
Sm __ .
_oC
\
\
7 matrix
DTA of heating (melting) (see initial mlcrostructure in Appendix AI)
and cooling (solidification) of Specimen i0 (CS).
BI.3.1
ORIGINAL PAGE IS
OF POOR QUALITY
|
U_u
k
|._
oll coo1_
_,mm mm.m m.m u_.m u_.oo ,,J,.m ,,,6.ram
TEMPERATURE (t:)
DTA of heating (melting) (see initial microstructure in Appendlx AI)
and cooling (solidification) of Specimen 27 (CS+).
BI.3.2
,_.';_,_,:,-',_. PAGE iS
OF POOR QUALITY
m |.m
0
on cooling
eating
t.|!w
_rt_
O.w_Q.oe
7 matrix
n m _ m _.m m_ m in'_R t_ m _t_ R t_ m
DTA
DTA of heating (melting) (see initial microstructure in Appendix AI)
and cooling (solidification) of Specimen 12 (BC).
BI.4
OE_GIr_AL PAGE IS
OF POOR QUALITY
on cooling
7 matrlx
f
m |.l
w0
III
J
on heating
m.m l_.m I_ m m_m n_ m
D'r&
DTA of heating (melting) (see initial mlcrostructure in Appendlx AI)
and cooling (solidification) of Specimen 13 (B).
BI.5
h,
L9luo
on cooling
7 matrix
on heating
7 matrix
o.o0_.am _ oo _ oo _.oo ,t_oo ._.u _oo _.u ioo
r',
DIA
DTA of heating (melting) (see initial microstructure in Appendix AI)
and cooling (solidification) of Specimen 20 (interstitial free).
BI.6
OF POOR QUALITY
m l.|
O
O_=
trt
= "_.= m_m.= uJ*.= u_*.= n.= ,,a,.uo ,-6.m --<W
TEMPERATURE [C) DTA
DTA of heating (melting) (see initial microstructure in Appendix AI)
and cooling (solidification) of Specimen 22 (C+).
BI.7
ORiG!NAL PAGE IS
OF POOR QUALITY
m l,l.ihiE0WJ
I.--Illl I0
on coolLng
Ik.= m.mo ld..= _,_*.= ,,-.= -_.m
TEMPERATURE (C) DTA
DTA of heating (melting) (see initial microstructure in Appendix AI)and cooling (solidification) of Specimen 25 (S+).
BI.8
OF POOR QUALITY,
7 matrix
Laves
f
In I,II,IIJIiJZmhJQ
O'II _.lO
7 matrix
Ioo.oo Ido oo oolTo.w oub.w t6t.oo o_ ai, aHo.oo omdo.oo nio.oo ,,-r,,.oo
DTA of heating (melting) (see initial mlcrostructure in Appendix AI)and cooling (solidification) of Specimen 14 (P).
BI.9
ORIGIN._L PAGE IS
OF POOR QUALITY
__ i'm I
20Fro
on cooling
7 matrix
NbC
Q.m "&*.m
on heating
NbC
C'-
DTA
DTA of heating (melting) (see initial microstructure in Appendix AI)and cooling (solidification) of Specimen 718.
BI.10
I.U
In 4.11UJLLIn.O
a
CAST 71111181 NF.LTZNG
_; I.N all _ NTU N.N IolBln
/LTIIBBq4EJt It).IUlll N L'I:SIDC,n
O,IIII .n _.N _.m _ .m IIII .R
B2 Reproducibility of DTA test data is excellent (±5C) as shown in 3runs of commercial alloy 718.
B2.1 ORiGIN_',L PAGE !$
OF POOR QUALITY
m LIDI/
!
IR 4.11LUt_
iWa
CJUIT 7111 |2| I_ NI[LTXNG
IT: l_l. |i ell _ llITl; U.N I_|,,
&qlWqiDlt Nil.lUre N Imli|m
iii
IJl.
I CAlll T|| (2J IOLTDIFTCAT|ON
li_lW,N[J_ Itl.|im N eeSoin
Lie
il.mtko
ia
.,.oo.oe 10oo/o|n
r
CAST Till |_) DSELT|NG
iii: 1141, II iql lCim RLLTIr:
J_rllD_lFJl_ PlEL|UIL I_ ¢¢/m1_
I1.tl Ill/lla
r-
I.II-L_too oo m,# _.i | .N llN .i II .I _ ,N t,i H.N
Commercial 718 alloy run number 2.
ORIGINAL PAGE IS
OF POOR QUALITY B2.2
I
LI
!
H Ill |11 II#,U II.N
Commercial 718 alloy run number 3.
B2.3OR_GIN_.L PAGE IS
OF POOR QUALITY
Appendix C
Hicrost__ct_res froR
Cleeble Weld Simulation
C
1210
1230 1250
Alloy 3: S. B. (P)
Microstructures after quenching from I190C, 1210C, 1230C, and 1250Cusing the Gleeble cycle of Figure I.
CI.1 ORtG_,_A'L FAGE IS
OF POOR (_UAI.i'rY_
1250
Alloy 5: C. S. B. (+)
Microstructures after quenching from 1190C, 1210C, 1230C,using the Gleeble cycle of Figure 1.
and 1250C
CI.2
ORIGINAL PAGE IS
OF POOR QUALITY
1190.i 1210
,. -._.. ._ _ ._ .-
I_'___ '._.llk-__-"_, ".
1250
Alloy I0; $. c
Microstructures after quenching from I190C, 1210C, 1230C,using the Gleeble cycle of Figure I.
and 1250C
Ci.3 ORICi|!_,_ALPAGE IS
OF POOR QUALITY
1190 1210
1230
Alloy 12: B, C
Microstructures after quenching from 1190C, 1210C, 1230C, and 1250Cusing the Gleeble cycle of Figure I.
CI.4ORIGINAL PAGE IS
OF POOR QUALITY
1190 1210
1230
Microstructures after quenching from I190C, 1210C, 1230C, and 1250Cusing the Gleeble cycle of Figure 1.
Ci.5OF POOR QUALITY
1190
Alloy 14:
Microstructures after quenching from i190C, 1210C, 1230C, and 1250Cusing the Gleeble cycle of Figure I.
i
CI.6ORIGINAL PAGE IS
OF POOR QUALITY
1190
Alloy 20; _nterstltial Free
Hicrostructures after quenching from i190C, 1210C, 1230C, and 1250C
using the Oleeble cycle of Figure 1.
Cl.7 ORIG!I_AL PACE _S
OF POOR QUALITY,
Alloy 21: S (+)
Mlcrostructures after quenching from 1235C using the Gleeble cycle ofFigure i.
CI.8ORiGiNAL PAGF. IS
OF POOR QUALITY
r__.-._ ......._:...
I_?:\-'I _ -___.:-. >,. -
...._:_ .... .
Alloy 22: C (+)
Mlcrostructures after quenching from 1215C and 1235C using the Gleeble
cycle of Figure i.
CI.9
Appendix D
Scanning Auger Microscopy (SAM)
D1.D2.
D2.1D2.2D2.3D2.3AD2.3BD2.4D2.5D2.6D2.7D2.8
Hydrogen KmbriCtlement ApparatusSelected SAM Microanaylsis Spectra and
Associated MicrogTaphsAlloy 3 (BSP)Alloy 5 (BSC+)Alloy 10 (CS)(1093C) Alloy 27 (CS+) Homogenized 1 hour at 1093C(927C) Alloy 27 (CS+) Homogenized 1 hour at 927CAlloy 12 (BC)Alloy 13 (B)Alloy 20 (Clean)Alloy 22 (C+)Alloy 25 (S+)
D
anode
cathode
basket support
II
sample support
platinum _ire
J
i
sample
2Jplatinum basket
stirrer
< >
magnetic stir-plate
i000 ml beaker
Charging Conditions
current density
sample surface area
electrolyte
electrolyte temperaturestir rate
charging time
= 400 ma/cmZ0.5 - 1.0 cm2
0.5 m H2SO_ + 50 mg NaAs02 per 1ambient
slow
13 - 21 days
Hydrogen charging cell used to embrlttle grain boundaries for exposure
in the scanning Auger microscope.
D1
File : 9005001AJ.SSP
lO0 ,
z 0 .,. ri tii
g: ,
20O 4OO 6OOKINETIC ENERGY (eV)
800 tO00
Alloy 3 (BS+P) Spot #50b8
SAM chemical spectra of a freshly fractured grain boundary surface.
Significant levels of B, S, and P are present. Note that C and 0 are
typical system contaminates.
D2.1.1
File : 900500tAL.SSP
tO00
1"4
_; * _b c _ Li _
v IIA)'
4/;Fe
-t000
iL
200 400 600 800 tO00KINETIC ENERGY (eV)
Alloy 3 (BS+P) Spot #50ci0
SAM chemlcal spectra
Significant levels of
typical contaminates.
of a freshly fractured matrix surface.
only S are present. Note that C and 0 are
D2.1.2
,_IG,_,L_L PAGE !S_| %= =.
Of poOR QUALI'T_
File : 900570tAH.SSP
t;:t-t
5OO
_p
-500
Cv
¢. ,-_r 1 [4fT..
.,,.1-- IV:Fc
]• • . _J
200 400 600 800KINETIC ENERGY (eV)
/J(
!1000
ON
Alloy 5 (BSC+) Spot #57c8
SAM chemical spectra of a freshly fractured grain boundary surface.
Significant levels of only P are present. Note that C and O are
typical system contaminates.
D2.2.1
ORIGIP_II PAGE IS
OF POOR QUALITY
File : 9005701BD.SSP
F r .......... ,Ii ...........400
, r Ii(.r £r i,
-20O
-400,
200 400 600 800 1000KINETIC ENERGY (eV)
Alloy 5 (BSC+) Spot #57c8
SAM chemical spectra of the grain boundary surface of D2.2.1 after
sputter cleaning. Significant levels of none of the dopants are
present. Note that C and 0 are typical system contaminates.
D2.2.1B
aoo
loo
i°-tO0
-20O
File : 900570tAS.SSP
r-,
rr
p
rV
Iv!
(/!
200 400 600 800 tO00KINETIC ENERGY (eV)
Alloy 5 (BSC+) Spot #57d15
SAM chemical spectra of a freshly fractured matrix surface.
Significant levels of none of the dopants are present. Note that C and
O are typical system contaminates.
D2.2.2
l-i
4OO
2OO
-20O
-400
Fill : 9005701AY.SSP
it_lI
!'l _
5
T"
[
200 400 600KINETIC ENERGY (eV)
800 lO00
Alloy 5 (BSC+) Spot #57e21
SAM chemical spectra of a freshly fractured carbide
Significant levels of only S are present. Note that 0 is
system contaminate.
surface.
a typical
D2.2.3
File : 9005701BF.SSP
2O0 rf
i
,oo"-
i °
-200200 400 600 BOO 1000
• K;NET_[C lr_;£ F4GY (eY)
Alloy 5 (BSC+) Spot #57e21
SAM chemical spectra of the carbide surface of D2.2.3 sputter cleaned
for 1 minute. Significant levels of only S are present. Note that O
is a typical system contaminate.
D2.2.3B
ORrGrNAL PAGE iS
OF POOR QUALI'r'_,
File : 9005701BJ.SSP
2OO
o
tOO 1 I
l!i I!l1|b 1"1 ,
-100 k_ . 1
-200
200.a J i I • e a t o t e I
400 600 800 1000KINETIC ENERGY (eV)
Alloy 5 (BSC+) Spot #57e21
SAM chemical spectra of the carbide surface of D2.2.3 sputtered cleaned
for 3 minutes. Significant levels of none of the dopants are present.
Note that 0 is a typical system contaminate.
OR;GINAL PAGE IS
OF POOR QUALITY
D2.2.3C
File : BOO850&AJ.BSP
i 0
-500
4_ G
ct
_00| • • • | •
40O 800KINETIC ENEREY (eV)
8O0 tO00
Alloy i0 (CS) Spot #85a7
SAM chemical spectra of a freshly fractured grain boundary surface.
Significant levels of S and B are present. Note that C and O are
typical system contaminates.
D2.3.1
Fi2e : 9008b01_5.S$_
kf
_o0
0
I-I
-200 -
,,,,..
c1
/
-..IL ..... m.. • .o I o 6 6 I 6 i • !
200 400 600 800K1;._.-1lC E;.;_.FIG'_ re', )
d(
• I J
1000
Alloy i0 (CS) Spot #85ai
SAM chemical spectra of freshly fractured grain boundary surface.
Significant levels of S and B are present. Note that C and 0 are
typical system contaminates.
D2.3.2
File : 900850|AS.SSP
f
200 400 600 800 iO00KINETIC ENERGY (eV)
Alloy I0 (CS) Spot #85c15
SAM chemical spectra of freshly fractured matrix surface. Significant
levels of only S are presen=. Note that C and 0 are typical system
contaminates.
D2.3.3
ORIGINAL PAGE ISOF pOOR QUALITY
Ftls : 9008501AT.SSP
tO0
>-F-t-4(nZbJI--Zi,.4 o
-iO0
200 400 600 800 |000KZNETZC ENERGY (eY)
Alloy 10(SO) Spot #85c16
SAM chemical spectra of a freshly fractured carbide interface surface.
Significant levels of only S are present. Note that C and 0 are
typical system contaminates.
D2.3.4
File : 90tOSOtAN.SSP
l-Iin
z
, I1 ,
li S lb c _i i Il> 17,1 .. il,i
Ii
O;
_.q. I;
ii
-2OO.f
200 400 600KINETIC Eltl_ft(;Y (eVl
#,'
80O lO00
Alloy 27 (CS+) as cast Spot #101elO
SAM chemical spectra of a freshly fractured grain boundary surface.
Significant levels of only S are present. Note that C and 0 are
typical system contaminates.
D2.3Acast.l
(.f
I4:
• ,a • i •
200 400 800 800 JO00KINETIC ENERGY (eV)
Alloy 27 (CS+) as cast Spot #101el0
SAM chemical spectra of the grain boundary surface after sputter
cleaning. Significant levels of only S (very small) are present. Note
that C and 0 are typical system contaminates.
D2.3.Acast. IB
I-4
-_00
-4OO
1";
C1F
200 400KINETIC ENERGY (eV)
6OO
/ I/vlill
1_V ,
. I'_av _
Fe
i
.a, • •
800 |000
Alloy 27 (CS+) as cast Spot #101c7
SAM chemical spectra of a freshly fractured matrix surface.
Significant levels of none of the dopants are present. Note that C and0 are typical system contaminates.
D2.3Acast.2
200
>-
I-4
o
M
-200
200 400 600 800 iO00KINETIC ENERGY (eV|
Alloy 27 (CS+) as cast Spot #101d9
SAM chemical spectra of a freshly fractured carbide surface.
Significant levels of only S are present. Note that C and 0 are
typical system contaminates.
D2.3Acast.3
>-k-M
b-4
8OO
0
-5OO
Flle: 9012701AE.SSP
=, .,
Cr
!
$
A
200 400 600 800KINETIC ENERGY (eV)
iO00
Alloy 27(CS+) Homogenized 1 hour at I093C Spot #127c5
SAM chemical spectra of a freshly fractured Sround boundary surface.Significant levels of only S are present. Note that C and 0 are
typical system contaminates.
D2.3B(1093).I
File : 90t270tBU.SSP
).-
k"l-tm
,.=,
I-I
5OO
0
-5OO
I
/it /Ft. r¢
C¢
6O0(ev)
r_
.°
r
20O 400 8O0 lO00KINETIC ENERGY
Alloy 27(CS+) Homogenized i hour at I093C Spot #12763
SAM chemical spectra of a sputtered grain boundary surface.
Significant levels of none of the dopants are present. Note that C and
O are typical system contaminates.
D2.3B(1093).2
Ftle: 90i2701AN.SSP
t-t
I-I
Imo
lO0
0
-iO0
-_00
ii : i
F,.
L=O0 400 800 BOO |000KINETIC ENERGY (eV)
Alloy 27(CS+) Homogenized 1 hour at I093C Spo_ 127d12
SAM chemical spectra of a cleaved carbide surface. Significant levels
of none of the dopants are present. Note that C and 0 are typicalsystem contaminates.
D2.3B(1093).3
File : 901270tBG.SSP
W
zuJ
I-4
5OO
I
i
tl0
Cv
-500
200 400 600 800 tO00KINETIC ENERGY (eV)
Alloy 27(CS+) Homogenized 1 hour at I093C Spot 127142
SAM chemical spectra of a freshly fractured carblde/matrix interface
surface. Significant levels of S and maybe B are present. Note that 0is a typical system contaminate.
D2.3B(1093).4
>-k-_4
Zp4
5OO
-50O
File : 90_2701AR.SSP
I
CT: Cr
2OO 400 6OOKZNETZC ENERGY (eV)
8O0 l0
Alloy 27(CS+) Homogenized 1 hour at 1093C Spot 127g17
SAM chemical spectra of a freshly fractured eutectlc surf_
Significant levels of only S are present. Note that C and 0
typical system contaminates.
D2.3.B(1093).5
500
0
k4
-5OO
Flle: 9004301AB.SSP
A , ,
\_ e"
ct
g',
g;
| • • L,
_00 400 600 800 1000KINETIC ENERGY (eV)
Alloy 12 (BC) Spot #43a2
SAM chemical spectra of freshly fractured grain boundary surface.Significant levels of only B are present. Note that C and 0 are
typical system contaminates.
D2.4.1
I
tO00 I
0 -' 1 I'
J/;
C¢-tO00 .....
L)O0 400 800 800 1000KIHETIC EHERGY (eY)
Alloy 12 (BC) Spot #43b3
SAM chemical spectra of freshly fractured grain boundary surface•
Significant levels of only B are present• Note that C and 0 are
typical system contaminates.
D2.4.2
500
Ft;e : 9004301AK,SSP
-500 ...... _"
JI
200 400 600 800 tO00KINETIC ENERGY (eY|
Alloy 12 (BC) Spot #43bi0
SAM chemical spectra of freshly fractured matrix surface. Significant
levels of none of the dopants are present. Note that C and 0 are
typical system contaminates.
D2.4.3
50O
File : 9003tOtAD.SSP
I
i
200 400 600 800 |000KINETIC ENERGY (eV)
Alloy 13 (B) Spot #31a4
SAM chemical spectra of a freshly fractured grain boundary surface.
Significant levels of only B are present. Note that C and 0 are
typical system contaminates.
_ D2.5.1
5OO
-5OO
File : 9003101AH.5_
I !
C," _/l
2OO 400 600 800KINETIC ENERGY (eV)
lO00
Alloy 13 (B) Spot #31b8
SAM chemical spectra of a freshly fractured matrix surface.Significant levels of none of the dopants are present. Note that C andO are typical system contaminates.
D2.5.2
File : 90t570tAC.SSP
-2OOMI
200 400 600 800 |000KINETIC ENERGY (eV)
Alloy 20 (clean) Spot #157a3
SAM chemical spectra of a freshly fractured grain boundary surface.
Significant levels of only B or C1 (contaminates) are present. Notethat C and 0 are typical system contaminates.
D2.6.1
C1"
-2OO
I!
2O0 400 600 BOOKINETIC ENERGY (eV}
I
fO00
Alloy 20 (clean) Spot #157a3
SAM chemical spectra of the grain boundary surface after sputter
cleaning. Significant levels of none of the dopants are present. Note
that C and O are typical system contaminates.
D2.6.1B
iO00
-lO00
RO0 400 500 BOO tOO0KINETIC ENERGY (eV)
Alloy 20 (clean) Spot #157c7
SAM chemical spectra of the grain boundary surface. Significant levels
of only B and C1 (contaminates) are present. Note that C and 0 are
typical system contaminates.
D2.6.2
200
i°-20O
File : 0015701_._
200 400 600 800KXNETXC ENER6Y (eV)
&,
lO00
Alloy 20 (clean) Spot #157d12
SAM chemical spectra of the matrix surface. Significant levels of only
B (small) are present. Note that C and 0 are typical systemcontaminates.
D2.6.3
I-4
aoo
-20o
File : 90i5701AP.SSP
iS
/
VD
r_
J
r
LJO0 4OO 6OO 8OOKINETIC ENERGY (eV)
7ii
SO00
Alloy 20 (clean) Spot #157f14
SAM chemical spectra of a freshly fractured precipitate zone
Significant levels of only S and B or C1 (contamination) are
Note that C and 0 are typical system contaminates.
surface.
present.
D2.6.4
File : 9017_OtAF.SSP
ItOO
i o - !t ge, :'
-lO0
_00 400 600 800 tO00KINETIC ENERGY (eV)
Alloy 22 (C+) Spot #171c6
SAM chemical spectra of a freshly fractured grain boundary surface.
Significant levels of S (small) and P (small) are present. Note that C
and O are typical system contaminates.
D2.7.1
Ftle: 9017101AS.SSP
I00 lJ
Fe _
_.o li
t
I !iCr
/9;_• . |
200 400 800 800 |000KINETIC ENERGY (eV)
Alloy 22 (C+)Spot #171c6
SAM chemical spectra of the grain boundary surface after sputter
cleaning. Significant levels of none of the dopants are present. Note
that C and 0 are typical system contaminates.
D2.7.1B
0=4
600
-5OO
File : 9017t0tAH.SSP
,., I
FI
I t #;
e,t
Nt
200 4O0 600 800KINETIC ENERGY (eV)
1000
Alloy 22 (C+) Spot #171d8
SAM chemical spectra of the grain boundary matrix surface.
levels of none of the dopants are present. Note that C
typical system contaminates.
Significantand 0 are
D2.7.2
File : 90i7iOIAJ.SSP
800 ,
¥;
L
-5OOaO0 400 800 800 SO00
KINETIC ENERGY (eV}
Alloy 22 (C+) Spot #171e9
SAM chemical spectra of a freshly fractured carbide surface.
Significant levels of none of the dopants are present. Note that C and0 are typical system contaminates.
I)2.7.3
Ftlo : IOfTfOf,td4.;rip
8o
200 400 1500KINETIC ENERSY (sV)
800 |000
-
Alloy 22 (C+) Spot #171g12
SAM chemical spectra of a freshly fractured eutectic surface.
Significant levels of only S are present. Note that C and 0 are
typical system contaminates.
D2.7.4
File : 90t6901AD.SSP
-400
L:rC!
i i
ZOO 400 80OKINETIC ENERGY (eV]
° .
800 lO00
!
Alloy 25 (S+) Spot #169b4
SAM chemical spectra of a freshly fractured grain boundary surface.
Significant levels of only B or Cl (contamination) are present. Note
that C and 0 are typical system contaminates.
D2.8.1
50O
0
_4
-5OO
Flle: 90f690tAV.SSP
I
, ,,,.
¢;AI_v
,L;/i!,i
Ii!J _c,, 1 T_
MJ
• , a
200 400 600 800 SO00KINETIC ENERGY (eV)
I
Alloy 25 (S+) Spot #169b4
SAM chemical spectra of the grain boundary surface after sputter
cleaning. Significant levels of none of the dopants are present. Note
that C and 0 are typical system contaminates.
DI. 8. IB
SO00
File : OO1600i,4L..8_ °
0.4
-|000- | - . .
800
I . . • i .
4O0 6OOKINETIC ENERGY (eV)
JN'• | f. .
800 tO00
Alloy 25 (S+) Spot #169d12
SAM chemical spectra of a freshly fractured matrix surface.
Significant levels of none of the dopants are present. Note that C and
0 are typical system contaminates.
D2.8.2
M
SO0
-SO0
Fl|o : 11016901AN.ESP
IC
1"
-r;
;lO0 400 600 000KINETIC ENERGY (or)
I •
|000
Alloy 25 (S+) Spot #169d13
SAM chemical spectra of a freshly fractured carbide surface.
Significant levels of only S are present. Note that C and 0 are
typical system contaminates.
D2.8.3
40O
@00
E_ o#=4
-@00
-4O0
Ftle : 9016901AE.SSP
A
)
i#
1-1
I
f
- a a, •
LIO0 400 BOO SO00KINETIC ENERGY (eV)
Alloy 25 (S+) Spot #169b5
SAM chemical spectra of a freshly fractured eutectic surface.
Significant levels of only S are present. Note that C and 0 aretypical system contaminates.
D2.8.4