nfsd 53-55 - nasa · 2013-08-30 · establish that sulfur and phosphorus do segregate to the grain...

NFSD 89-0 (June 30, 1989)

FORMS

53-55

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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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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 )

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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


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


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


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


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


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


Significant levels of only P are present. Note that C and O are


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!

(/!


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)


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



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


Significant levels of S and B are present. Note that C and O are


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


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


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




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


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.



D2.3Acast.3

>-k-M

b-4

8OO

0

-5OO

Flle: 9012701AE.SSP

=, .,

Cr

!

$

A


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


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.


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


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


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


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


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


D2.4.3

50O

File : 9003tOtAD.SSP

I

i

200 400 600 800 |000KINETIC ENERGY (eV)

Alloy 13 (B) Spot #31a4


Significant levels of only B are present. Note that C and 0 are


_ 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


Alloy 20 (clean) Spot #157a3


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


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


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


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;_• . |


Alloy 22 (C+)Spot #171c6




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


Significantand 0 are

D2.7.2

File : 90i7iOIAJ.SSP

800 ,

¥;

L

-5OOaO0 400 800 800 SO00

KINETIC ENERGY (eV}

Alloy 22 (C+) Spot #171e9


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.



D2.7.4

File : 90t6901AD.SSP

-400

L:rC!

i i

ZOO 400 80OKINETIC ENERGY (eV]

° .

800 lO00

!

Alloy 25 (S+) Spot #169b4


Significant levels of only B or Cl (contamination) are present. Note


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





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



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




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)


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

nfsd 53-55 - nasa · 2013-08-30 · establish that sulfur and phosphorus do segregate to the grain...

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