Atlas of Microstructures

iAtlas of Microstructures

Atlas of MicrostructuresAF234/7226/MSchü

Prepared by: Ellen Berghof-Hasselbächer 1)

Peter Gawenda 1) Monika Schorr 1) Michael Schütze 1) John J. Hoff man 2)

1) Karl-Winnacker-Institut der DECHEMA e.V.Th eodor-Heuss-Allee 25

60486 Frankfurt am MainGermany

2) Air Products and Chemicals, Inc.7201 Hamilton BoulevardAllentown, PA 18195-1501

USA

Th e results, conclusions and recommendations given in this report refer to the specimens and data submitted as well as to the system conditions mentioned. Th erefore, the results cannot be applied to other conditions. DECHEMA e.V. accepts no liability.

It is only permitted to reproduce this report unabridgedly. Publication of the report or of single experimental results are subject to approval by DECHEMA e.V.

ii Atlas of Microstructures

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All rights reserved, including translations

ISBN 978-1-57698-046-0

No part of this publication may be reproduced, stored in a retrieval system, or transmitted, in any form or by any means, electronic, mechanical, photocopying, recording or otherwise, without prior written permission of the publisher. This document was prepared under the sponsorship of the Materials Technology Institute Inc. (MTI) and is approved for release. All data and information contained in this document are believed to be reliable; however, no warranty of any kind, express or implied, is made with respect to the data, analyses, or author of this document; and the use of any part of the document is at the user's sole risk. MTI, the author, or any person acting on its behalf, assume no liability and expressly disclaim liability, including without limitation liability for negligence, resulting from the use or publication of the information contained in this document to warrant that such use or publication will be free from privately owned rights.

Published by Materials Technology Institutewww.mti-global.org

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Table of Contents

Introduction ........................................................................................................................... xx Description of Test Specimens ................................................................................................ xxCreep rupture test specimens........................................................................................................ xx

- Alloy 35/45 Exposure Conditions ..................................................... Table 1- Chemistry of Alloy 35/45 Creep Rupture Stress Samples .................. Table 2- Alloy HPMA Exposure Conditions .................................................. Table 3- Chemistry of Alloy HPMA Creep Rupture Stress Samples ................ Table 4- Alloy HP Modified Exposure Conditions ......................................... Table 5- Chemistry of Alloy HP Modified Creep Rupture Stress Samples ....... Table 6

Service-exposed specimens ........................................................................................................... xx- History of Service Exposed Samples- Exposure Conditions of ................................................................... Table 7- Chemistry of Service Exposed Samples .............................................. Table 8

Sample Preparation and Microstructural Characterization ...................................................... xxMetallographic Preparation ..................................................................................................... xxInvestigation Techniques ............................................................................................................. xx

- Differential Interference Contrast Technique (LM-DIC)- Scanning Electron Microscopy (SEM-BSE)- Electron Probe Microanalysis (EPMA)- Interference Layer Metallography (LM-ZnSe)- Image Analysis by the False Color Technique (LM-FC)- Development of the Microstructure by Etching (LM-etched)

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Results of the Microstructural Analyses ................................................................................... xxAlloy 35/45 ............................................................................................................................... xx

- Results of image analysis of Alloy 35/45 ............................................ Table 9- As Cast .............................................................................................. Fig. 1-1a-j- Creep Rupture Samples .................................................................... Fig. 1-2a to 1-2j .................................................................... Fig. 1-3a to 1-3r .................................................................... Fig. 1-4a to 1-4r .................................................................... Fig. 1-5a to 1-5r .................................................................... Fig. 1-6a to 1-6z .................................................................... Fig. 1-7a to 1-7r .................................................................... Fig. 1-8a to 1-8r .................................................................... Fig. 1-9a to 1-9z .................................................................... Fig. 1-10a to 1-10r .................................................................... Fig. 1-11a to 1-11r .................................................................... Fig. 1-12a to 1-12z .................................................................... Fig. 1-13a to 1-13v- Conclusions ......................................................................................................................... xx

Alloy HPMA ............................................................................................................................. xx- Results of image analysis of Alloy HPMA .............................Table 10- As Cast .............................................................................................. Fig: 2-1a to 2-1j- Creep Rupture Samples .................................................................... Fig. 2-2a to 2-2r .................................................................... Fig. 2-3a to 2-3r .................................................................... Fig. 2-4a to 2-4r

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- Creep Rupture Samples (continued) ...................................................... Fig. 2-5a to 2-5r .................................................................... Fig. 2-6a to 2-6r .................................................................... Fig. 2-7a to 2-7z .................................................................... Fig. 2-8a to 2-8r .................................................................... Fig. 2-9a to 2-9r .................................................................... Fig. 2-10a to 2-10r .................................................................... Fig. 2-11a to 2-11r .................................................................... Fig. 2-12a to 2-12r .................................................................... Fig. 2-13a to 2-13r .................................................................... Fig. 2-14a to 2-14r .................................................................... Fig. 2-15a to 2-15r .................................................................... Fig. 2-16a to 2-16r .................................................................... Fig. 2-17a to 2-17s- Conclusions ......................................................................................................................... xx

Alloy HP Modified ..................................................................................................................... xx- Results of image analysis of Alloy HP Modified ................................. Table 11- As Cast ............................................................................................. Fig. 3-1a to 3-1j- Creep Rupture Samples .................................................................... Fig. 3-2a to 3-2r .................................................................... Fig. 3-2a to 3-2r .................................................................... Fig. 3-3a to 3-3r .................................................................... Fig. 3-4a to 3-4r .................................................................... Fig. 3-5a to 3-5r .................................................................... Fig. 3-6a to 3-6r .................................................................... Fig. 3-7a to 3-7r

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- Creep Rupture Samples (continued) ...................................................... Fig. 3-8a to 3-8r .................................................................... Fig. 3-9a to 3-9r .................................................................... Fig. 3-10a to 3-10r .................................................................... Fig. 3-11a to 3-11r .................................................................... Fig. 3-12a to 3-12r .................................................................... Fig. 3-13a to 3-13r- Conclusions ......................................................................................................................... xx

Service Exposed Samples ............................................................................................................. xx- Results of image analysis of Service Exposed Samples ........................ Table 12- Alloy HPMA ................................................................................... Fig. 2-SE-Fa to 2-SE-Fr .................................................................... Fig. 2-SE-Ga to 2-SE-Gr- Alloy HP Modified .......................................................................... Fig. 3-SE-Ba to 3-SE-Br .................................................................... Fig. 3-SE-Aa to 3-SE-Ar .................................................................... Fig. 3-SE-Ca to 3-SE-Cr .................................................................... Fig. 3-SE-Da to 3-SE-Dr .................................................................... Fig. 3-SE-Ha to 3-SE-Hr .................................................................... Fig. 3-SE-Ia to 3-SE-Ir .................................................................... Fig. 3-SE-Ja to 3-SE-Jr .................................................................... Fig. 3-SE-Ea to 3-SE-Ez- Conclusions ......................................................................................................................... xx

Summary ................................................................................................................................ xxComposition of the Precipitates .................................................................................................... xx

Average Compositions of the Precipitates in atomic-% ......................... Table 13

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Precipitation Kinetics Diagrams .............................................................................................. xx- Alloy 35/45 ..................................................................................... Fig. 1-PKD-a to 1-PKD-d .................................................................... Fig. 1-PKD-e

- Alloy HPMA .................................................................................. Fig. 2-PKD-a to 2-PKD-c .................................................................... Fig. 2-PKD-d

- Alloy HP Modified ................................................................... Fig. 3-PKD-a to 3-PKD-d .................................................................... Fig. 3-PKD-e

�Atlas of Microstructures

IntroductionUp to and through the 1940’s, furnace tubes were typically fabricated

from wrought chromium steels and/or austenitic stainless steel alloys. Having rather low carbon concentrations, these alloys had poor creep strength and limited service lives. To increase tube life, greater carbon concentrations were required to promote precipitation hardening upon elevated temperature exposure. The increased carbon concentrations, however, generally resulted in excessive work hardening during conven-tional processing of wrought materials. The use of centrifugal casting, pouring molten metal in a horizontal spinning mold, allowed founders to develop high carbon alloys as the molten metal solidified into near final shape without the need for subsequent metal working operations. Thus, refinement of centrifugal casting processes was the gateway to further alloy development and more aggressive furnace operations.

The first widely used centrifugal cast alloy for steam-methane reformer tube applications was HK40 in the 1950’s timeframe. HK40 is essentially the cast equivalent to wrought 310 stainless steel nominally containing 25 wt% chromium, 20 wt% nickel, with iron as the balance. However, HK40 nominally contains 0.40 wt% carbon while wrought 310 stainless steel contains only 0.08 wt% carbon. The increased carbon content and precipitation of primary carbides resulted in HK40 having greatly improved high temperature strength as compared to wrought 310 stainless steel. In the 1960’s, the cast HP alloys (nominally 25 wt% chro-mium, 35 wt% nickel, 0.50 wt% carbon, with iron as the balance) were developed to provide greater creep strength as compared to HK40.

The HK and HP alloys rely on precipitation of M23C6 and/or M7C3 carbides (where M is primarily chromium) for elevated temperature creep strength. The precipitated chromium carbides in the HK and HP alloys tended to coalesce as exposure temperatures approached 1800oF (982oC). The carbide coalescence decreased the creep strength and, therefore, limited the strength of these alloys at elevated temperatures. Microstructural changes that occur in the HK and HP alloys with

extended aging time and temperature have been well documented by Battelle Columbus Laboratories1. The Battelle data has proven to be a valuable resource in estimating reformer tube exposure conditions associ-ated with reformer tube failure analyses and remaining life assessments.

User demand for higher temperature/stronger alloys fueled continued alloy development resulting in the introduction of the HP-modified alloy in the 1970’s. The HP-modified alloy had the same nominal chemistry of the HP alloy along with the addition of typically 1 wt% niobium. The niobium addition results in precipitation of M23C6, M7C3, and MC type primary carbides upon solidification. In the M23C6 and M7C3 carbides, niobium substitutes for some of the chromium with the complex niobium-chromium carbides being more stable at elevated temperatures as compared to chromium carbides. In the HP-modified alloy, niobium is the primary carbide forming element in the MC type carbides.

In the 1980’s, the demand for more severe design conditions and/or design lives in excess of 100,000 hours led to the introduction of the HP-micro-alloyed material. HP-micro-alloyed (or commonly desig-nated as HPMA) material was based on the HP-modified chemistry with “micro” additions of alloying elements. In general, micro-alloying refers to intentional alloying additions at a concentration of 0.10 wt% or less. Titanium is the most common micro-alloying addition with other micro-alloying additions including zirconium, tantalum, or tungsten. Rare earth elements such as lanthanum, cerium, and/or yttrium may also be added. The micro-alloying additions provide a fine dispersion of MC type carbides that are stable at temperatures well in excess of 2,000oF (1093oC).

In the 1990’s, the 35Cr/45Ni alloy family grew in popularity. The 35Cr/45Ni alloy has similar creep strength to the HPMA alloy but with notably improved carburization resistance. The excellent carburization resistance made the 35Cr/45Ni alloy well suited for ethylene pyrolysis furnace tubes.

As outlined above, alloy development has continued in centrifugally cast heat resistant alloys allowing users to design for and operate at more

� Atlas of Microstructures

severe conditions. Unlike the HK40 and HP50 alloys, there is virtu-ally no published data on the microstructural changes that occur in the HP-modified, HPMA, or 35Cr/45Ni alloys upon long term aging. In an effort to address industry need, the Materials Technology Institute (MTI) sponsored the Atlas of Microstructures project. The specimens analyzed in the Atlas of Microstructures were foundry stress rupture specimens generously donated by MetalTek International, Duraloy Technologies, Kubota Metal Corp., and Manoir Industries. In addition, service exposed samples, having longer aging times than the foundry stress rupture specimens, were donated by operating companies includ-ing Air Products & Chemicals, Syncrude Canada, Eastman Chemical, and MetalTek International. The thorough, detailed microstructural analyses were completed at the Karl-Winnacker Institut der DECHEMA in Frankfurt, Germany.

The Atlas of Microstructures documents microstructural changes with increased aging time and temperature, identification and chemical composition of precipitated phases, as well as diagrams characterizing the kinetics of phase transformation for the HP-modified, HPMA, and 35Cr/45Ni alloy classes. This MTI Atlas of Microstructures bridges the gap from the Battelle project from the early 1970’s to the most impor-tant alloys used in the petrochemical industry today.

continued on page 8

Description of the Test SpecimensCreep rupture test specimens

The specimens that had been taken from creep rupture tests can be summarized with regards to their exposure conditions and their chemistry in Tables 1-6. For each of the three materials investigated an as-delivered (as-cast) specimen as reference sample was available. For alloy 35Cr/45Ni and alloy HP Modified twelve creep specimens had been supplied for the investigations. For alloy HPMA sixteen creep spec-imens existed. The test temperatures varied between 1675 and 2100°F (913 and 1149°C), the maximum exposure times reached 12,289 h. The stress range was between 1.0 and 8.5 ksi (6.9 and 58.8 MPa). The scatter of the chemical compositions was in the allowed range for these materials.

3Atlas of Microstructures

Table 1: Alloy 35/45 Exposure ConditionsSample_ID Dechema Sample MTI Test-Temp °F Test-Temp °C Stress Ksi Stress MPa Expo. Time hrs

A234_003 35/45-1 (c) as cast as castA234_004 35/45-2 1800 982 3.30 22.8 1,509A234_005 35/45-3 1800 982 2.90 20.0 7,217A234_006 35/45-4 1922 1050 3.40 23.4 367A234_007 35/45-5 1900 1038 2.09 14.4 2,023A234_008 35/45-6 1900 1038 1.60 11.0 6,331A234_009 35/45-7 1900 1038 1.54 10.6 10,247A234_010 35/45-8 2012 1100 2.00 13.8 335A234_011 35/45-9 2000 1093 1.80 12.4 1,297A234_012 35/45-10 2012 1100 1.15 7.9 3,545A234_013 35/45-11 2057 1125 1.60 11.1 277A234_014 35/45-12 2050 1121 1.00 6.9 2,606A234_015 35/45-13 2100 1149 1.50 10.3 248

Table 2: Chemistry of Alloy 35Cr/45Ni Stress Rupture SamplesSample_ID Sample Composition (wt%)Dechema MTI Cr Ni Fe Mn Si C Nb N OtherA234_003 35/45-1 (c) 35.85 44.59 bal 0.60 1.55 0.38 0.91 0.080 Ti, Zr & WA234_004 35/45-2 35.12 43.69 bal 1.32 1.83 0.46 0.78 0.060 Ti, Zr & WA234_005 35/45-3 34.84 44.43 bal 1.46 1.86 0.43 0.83 0.040 Ti, Zr & WA234_006 35/45-4 32.30 43.40 bal 1.02 1.52 0.41 1.21 0.065 TiA234_007 35/45-5 32.39 44.77 bal 1.40 1.85 0.42 1.25 NRA234_008 35/45-6 34.34 45.42 bal 1.04 1.30 0.37 1.22 NR TiA234_009 35/45-7 34.41 45.37 bal 1.31 1.78 0.42 1.20 NR TiA234_010 35/45-8 32.30 43.40 bal 1.02 1.52 0.41 1.21 0.065 TiA234_011 35/45-9 32.65 44.61 bal 1.37 1.71 0.42 1.14 NRA234_012 35/45-10 34.84 44.43 bal 1.46 1.86 0.43 0.83 0.040 Ti, Zr & WA234_013 35/45-11 32.30 43.40 bal 1.02 1.52 0.41 1.21 0.065 TiA234_014 35/45-12 32.59 44.34 bal 1.43 1.82 0.42 1.25 NRA234_015 35/45-13 34.84 44.43 bal 1.46 1.86 0.43 0.83 0.040 Ti, Zr & W

4 Atlas of Microstructures

Table 3: Alloy HPMA Exposure ConditionsSample_ID Dechema Sample MTI Test-Temp. °F Test-Temp. °C Stress Ksi Stress MPa Expo. Time hrs

A234_018 HPMA-1 (c) as cast as castA234_019 HPMA-2 1700 927 8.00 55.2 183A234_020 HPMA-3 1700 927 6.00 41.3 2,338A234_021 HPMA-4 1700 927 5.30 36.6 6,072A234_022 HPMA-5 1750 954 4.45 34.1 8,359A234_023 HPMA-6 1800 982 5.97 41.2 177A234_024 HPMA-7 1800 982 3.51 24.2 2,436A234_025 HPMA-8 1800 982 3.90 26.9 6,478A234_026 HPMA-9 1832 1000 3.00 20.7 12,289A234_027 HPMA-10 1922 1050 3.48 24.0 137A234_028 HPMA-11 1950 1066 2.80 19.1 2,558A234_029 HPMA-12 1922 1050 2.32 16.0 4,864A234_030 HPMA-13 1922 1050 2.08 14.3 11,778A234_031 HPMA-14 2012 1100 2.39 16.5 105A234_032 HPMA-15 2012 1100 1.80 12.4 2,714A234_033 HPMA-16 2012 1100 1.45 10.0 5,580A234_034 HPMA-17 2012 1100 1.35 9.3 8,990


Table 4: Chemistry of Alloy HPMA Stress Rupture SamplesSample_ID Sample Composition (wt%)Dechema MTI Cr Ni Fe Mn Si C Nb N OtherA234_018 HPMA-1 (c) 25.01 33.79 bal 0.65 1.08 0.45 0.65 0.06 Ti, Zr & WA234_019 HPMA-2 24.92 34,54 bal 0.78 0.80 0.50 0.49 NR TiA234_020 HPMA-3 24.29 32.51 bal 1.31 1.37 0.42 0.40 NR WA234_021 HPMA-4 25.01 33.79 bal 0.65 1.08 0.45 0.65 0.06 Ti, Zr & WA234_022 HPMA-5 25.01 33.79 bal 0.65 1.08 0.45 0.65 0.06 Ti, Zr & WA234_023 HPMA-6 25.85 35.25 bal 0.89 1.12 0.47 0.95 NRA234_024 HPMA-7 24.38 35.13 bal 0.79 0.85 0.46 1.08 NR Ti & WA234_025 HPMA-8 24.49 33.53 bal 0.56 1.14 0.44 0.51 0.05 Ti, Zr & WA234_026 HPMA-9 24.49 33.53 bal 0.56 1.14 0.44 0.51 0.05 Ti, Zr & WA234_027 HPMA-10 24.36 33.18 bal 1.37 1.73 0.46 0.83 NR Ti & ZrA234_028 HPMA-11 25.11 36.13 bal 0.91 1.33 0.48 NR Ti & WA234_029 HPMA-12 25.01 33.79 bal 0.65 1.08 0.45 0.65 0.06 Ti, Zr & WA234_030 HPMA-13 24.49 33.53 bal 0.56 1.14 0.44 0.51 0.05 Ti, Zr & WA234_031 HPMA-14 24.36 33.18 bal 1.37 1.73 0.46 0.83 NR Ti & ZrA234_032 HPMA-15 24.49 33.53 bal 0.56 1.14 0.44 0.51 0.05 Ti, Zr & WA234_033 HPMA-16 24.49 33.53 bal 0.56 1.14 0.44 0.51 0.05 Ti, Zr & WA234_034 HPMA-17 25.01 33.79 bal 0.65 1.08 0.45 0.65 0.06 Ti, Zr & W


Table 5: Alloy HP Modified Exposure ConditionsSample_ID Dechema Sample MTI Test-Temp. °F Test-Temp. °C Stress Ksi Stress MPa Expo. Time hrs

A234_036 HPNb-1 (b) as cast as castA234_038 HPNb-2 1675 913 6.50 44.6 659A234_039 HPNb-3 1700 927 8.50 58.8 59A234_040 HPNb-4 1750 955 5.10 35.2 794A234_041 HPNb-5 1778 970 5.15 25.5 286A234_042 HPNb-6 1778 970 3.77 26.0 1,185A234_043 HPNb-7 1800 982 5.40 37.4 191A234_044 HPNb-8 1800 982 5.42 37.4 4,467A234_045 HPNb-9 1800 982 5.52 37.4 7,833A234_046 HPNb-10 1800 982 2.80 19.3 10,637A234_047 HPNb-11 1850 1010 3.60 24.8 707A234_048 HPNb-12 1900 1038 2.30 15.8 2,555A234_049 HPNb-13 1900 1038 1.90 13.1 5,373


Table 6: Chemistry of Alloy HP Modified Stress Rupture SamplesSample_ID Sample Composition (wt%)Dechema MTI Cr Ni Fe Mn Si C Nb N OtherA234_036 HPNb-1 (b) 25.30 34.93 bal 1.29 1.40 0.44 1.06 NRA234_038 HPNb-2 25.34 34.88 bal 0.74 1.67 0.43 0.75 0.06 W & ZrA234_039 HPNb-3 25.26 33.56 bal 0.72 1.23 0.42 0.74 0.06 W & ZrA234_040 HPNb-4 25.34 34.88 bal 0.74 1.67 0.43 0.75 0.06 W & ZrA234_041 HPNb-5 25.11 32.85 bal 1.34 1.79 0.45 0.83 NRA234_042 HPNb-6 25.11 32.85 bal 1.34 1.79 0.45 0.83 NRA234_043 HPNb-7 24.88 33.33 bal 0.84 1.29 0.41 1.18 NR WA234_044 HPNb-8 24.86 33.79 bal 0.90 1.37 0.41 1.22 NR WA234_045 HPNb-9 24.86 33.30 bal 0.83 1.35 0.42 1.20 NR WA234_046 HPNb-10 25.26 33.56 bal 0.72 1.23 0.42 0.74 0.06 W & ZrA234_047 HPNb-11 25.62 34.75 bal 0.94 1.71 0.48 0.82 0.06 WA234_048 HPNb-12 25.26 33.56 bal 0.72 1.23 0.42 0.74 0.06 W & ZrA234_049 HPNb-13 25.26 33.56 bal 0.72 1.23 0.42 0.74 0.06 W & Zr


Metallographic Preparation

From the specimens available sections were taken for metallographic preparation with the surfaces to be investigated oriented parallel to the longitudinal direction of the tensile creep rupture specimens or perpen-dicular to the longitudinal axis of the tube specimens. Cutting was performed by using a precision sectioning machine with direct water cooling of the specimen. Some of the specimens had been delivered in the embedded state from which smaller sections were taken by saw cutting which were embedded again. Hot embedding in a conductive epoxy resin was used with a diameter of the moulds of 25 mm and after this the specimens were ground on SiC papers of the grades 180, 220, 320, 500, 1000 down to grit 2400 (Struers Standard 43-GB-1984, DIN 69176, Part 1,2,4) with a pressure of 70-80 N and water as a lubricant. Fine polishing was performed with a two step diamond polish of 3 μm and 1 μm followed by the finalizing polish with SiO2 suspension of size 0.02 μm. The area on the specimen surface to be investigated was marked by four Vickers hardness imprints (HV1) with a distance of 500 μm. Usually these areas were marked in the center of the specimen (in some cases, however, in addition regions close to the surface were investigated after marking).

Investigation TechniquesDifferential Interference Contrast Technique (LM-DIC)

For documentation of the microstructure in the light microscope at high magnification (500x and 1000x) the differential interference contrast (LM-DIC) technique2 was used by which the different phases due to the different hardness reveal certain topographical structures which facilitates distinction of these phases.

Scanning Electron Microscopy (SEM-BSE)After the LM-DIC investigations the same spot was investigated

by the scanning electron microscope using the back-scattered electron imaging technique (SEM-BSE). In this case the differences in contrast of the different precipitates resulting from differences in the density can be used in order to distinguish between the phases.

Electron Probe Microanalysis (EPMA)As a next step the phases documented by the other two techniques

were analyzed quantitatively for selected specimens with regard to their chemical composition. In each case several spot measurements were taken with a beam diameter of 2 μm. The results of these measure-ments were averaged. From the ratio between the different metals in the precipitates conclusions were drawn on the respective carbide type by taking carbon and nitrogen respectively from the difference of the sum of the determined metal fractions and 100%.

Interference Layer Metallography (LM-ZnSe)After the EPMA investigations the specimens were coated by a PVD

process in an evaporation equipment (Edwards) at a vacuum of about 10-3 mbar3-7. As a coating material zinc selenide (ZnSe) was used which allowed the comparison of these marked areas in the colored state with the images from the other techniques. The evaporation source consists of a little vessel made of tantalum with ZnSe grains filled into this vessel. By resistance heating the vessel is heated to a point where the ZnSe starts to evaporate. The evaporation rate is controlled by controlling the heating current together with observation of changes of the color of the section surface. When reaching the desired color (violet as a macro-scopic color) of the specimen surface the heating current is switched off immediately in order to achieve a reproducible thickness of the interfer-ence layer. Since the reflection characteristics change periodically with the layer thickness and manual control of the heating current is not easy this coating technique requires significant experience in order to achieve


a suitable interference layer. In some cases the layer had to be removed again from the section surface several times and a new coating had to be applied in order to come to satisfactory results. Due to the small dimen-sions of the different phases there may be slight deviations between the different photographs taken by the different techniques.

Image Analysis by the False Color Technique (LM-FC)By the use of the automated image analysis system Leica QWin in

combination with an automated laboratory microscope Leica DMLA the different fractions of the precipitations which had been characterized before by the other techniques were measured. For these measurements specimens were used which had been contrasted by the interference layer technique beforehand. An automated program routine was devel-oped by which for each specimen several focused images were taken at representative spots in the specimen center and partially also at the specimen edges at a magnification of 1000x. The images were stored and in a second run the lower and upper threshold values for the RGB (red, green, blue)-colors of the different phases were determined and recorded. Since, during the coating process the ZnSe layer can vary from specimen to specimen, this procedure had to be performed for all specimens at least once. For specimens where the microstructure in the center and in the outer region was different this procedure had to be applied separately for the two regions. The program routine allows manual interaction with the measurement procedure, i.e. artifacts like creep pores or cracks can be eliminated in order not to influence the measurement results. Based on the information in the binary memory, each of the phases was represented by a defined (false) color. For the different types of carbides and other phases the following false colors were selected: M7C3 - purple M6C - red M23C6 - yellow M2(C,N) - green M(C,N), MN - blue, cyan G-phase - magenta

The matrix was not binarized and was used as a background in true colors behind the false color image.

Development of the Microstructure by Etching (LM-etched)Finally the interference layers were polished off and sections were

etched for 30 seconds at 50°C in etchant V2A (composition: 100 ml H2O, 100 ml HCl 1.19, 10 ml HNO3 1.40, 0.3 ml Dr. Vogel’s pickle. Dr. Vogel’s pickle is a mixture of organic solvents with Thiourea. It consists of 1-Methoxy- 2-propanol (40-50%), Thiourea (3-5%) and Nonylphenol-ethoxylate (5-7%).

The interesting areas of the specimen were photographed at 200x and 500x magnification. By etching the matrix is partially removed so that edges are formed at the transition from matrix to precipitate. Due to the local reflection situation of the light which hits the surface perpendicularly the edges of these phase boundaries appear dark in the photograph so that all phases have a dark seam and the grain boundaries become visible.

(For more discussion on the etching technique, see the Addendum).

�0 Atlas of Microstructures

Table 9: Results of image analysis of Alloy 35Cr45Ni Creep Rupture Stress SamplesSample_ID Sample Temp. Temp °F Stress Stress Expo. area fraction in %

Time core outer zone / edge°C °F MPa Ksi hrs M(C,N) M�C3 M�3C� M�C M�(C,N) G-Phase M(C,N) M�C3 M�3C� M�C M�(C,N)

A234_003 35/45-1 (c) as cast as cast 0.8 3.9A234_004 35/45-2 982 1800 22.8 3.30 1,509 0.5 7.9 1.5 0.8 5.9 6.3A234_005 35/45-3 982 1800 20.0 2.90 7,217 0.3 7.8 1.8 0.3 6.6 12.5 +A234_006 35/45-4 1050 1922 23.4 3.40 367 0.4 6.9 0.7 0.3 7.4 4.2 +A234_007 35/45-5 1038 1900 14.4 2.090 2,023 < 0.05 8.3 < 0.03 3.0 < 0.02 8.8 13.1 +A234_008 35/45-6 1038 1900 11.0 1.60 6,331 < 0.05 10.5 7.5 < 0.05 + +++ ++A234_009 35/45-7 1038 1900 10.6 1.54 10,247 0.1 15.1 10.7 + +++ ++A234_010 35/45-8 1100 2012 13.8 2.00 335 1.1 7.6 1.1 7.8 6.8 +A234_011 35/45-9 1093 2000 12.4 1.80 1,297 0.4 10.8 0.3 1.2 0.2 10.2 10.8 ++A234_012 35/45-10 1100 2012 7.9 1.15 3,545 0.4 12.0 3.8 0.6 8.9 8.7 ++A234_013 35/45-11 1125 2057 11.1 1.60 277 1.1 7.2 0.2 7 4.5 +A234_014 35/45-12 1121 2050 .06.9 1.00 2,606 0.2 12.7 9.2 +++ ++A234_015 35/45-13 1149 2100 10.3 1.50 248 0.6 8.6 <0.3 0.6 9.9 7.6 ++

Annotations: + few, ++ many, +++ very many

��Atlas of Microstructures

Results of the Microstructural AnalysesAlloy 35Cr45Ni

The results from the image-analytical determination of the frac-tions of the precipitates in the microstructure are summarized in Table 9 (previous page) for all specimens investigated.

As-CastIn the as-cast condition the microstructure consists of an austen-

itic matrix with a eutectic mixture of carbides of the types M7C3 and M(C,N) which is situated between the dendrites and along the grain boundaries respectively. The fraction of M7C3 amounts to about 4%, while that of M(C,N) lies at about 0.8%. The M7C3 carbides are rich in chromium, while the M(C,N) carbides are rich in niobium and/or titanium, Figures 1-1a to 1-1j. In the DIC representation the M7C3

carbides appear light gray and the M(C,N) have a reddish appearance and are slightly elevated above the matrix due to their higher hardness, Figures 1-1a and 1-1b. In the BSE image the Nb-rich M(C,N) particles appear significantly brighter than the Cr-rich M7C3 particles and the Ti-rich M(C,N) which are almost black, Figures 1-1e and 1-f. After coating the specimen surface with ZnSe the MC particles containing high amounts of Nb appear in blue color, Figures 1-1g and 1-h, while the Ti-rich M(C,N) which could be detected in the DIC image by their strong intrinsic orange color were colored in light blue to white. The austenitic matrix and the M7C3 particles both have a reddish color, Figures 1-1g and 1-1h; however, the carbides can be distinguished quite clearly from the matrix as they are significantly darker. After etching the carbides, which have a dark seam, can be clearly distinguished from the matrix, Figure 1-1i. In this figure it becomes evident that these have been precipitated in regions of interdendritic residual melts or at grain boundaries, respectively. In the false color image the M7C3 particles were

colored purple and the M(C,N) were colored blue, Figure 1-1j.

Creep rupture samples

1800°F/982°C

��00°F/�,50� h/3.3 ksi (��°C/�50� h/��.� MPa)During exposure a transformation of the M7C3 particles had taken

place. The microstructure reveals some single creep pores, Figure 1-2a, and a dominance of M23C6 carbides (7.9%) which had also formed very fine secondary precipitates, Figures 1-2a to 1-2d. Furthermore a small amount (1.5%) of G-phase can be observed which is frequently in direct contact with the M23C6 carbides and shows a gray color with a slightly reddish tone in the DIC image. It cannot be excluded that among those particles which had been identified as G-phase and which have a dot-like appearance (i.e. very small particles) could also be M6C carbides in some spots. Due to their small dimensions it had not been possible to iden-tify these precipitates exactly either by their color contrast or by EPMA analysis. The G-phase (in the present analysis Ni16Nb6Si7) (according to the ASM Handbook of Metals also Ni18Ti8Si16

8) is supposed to have a negative influence on creep lifetime9. Due to its high amount of Nb this phase appears bright in the BSE image, Figures 1-2e and 1-2f. After coating with ZnSe the color becomes dark blue, while the Ti and Nb-rich M(C,N) particles (0.5%) which had appeared as orange precipitates in the DIC image become light blue to white, Figures 1-2c, 1-2d, 1- 2g and 1-2h. After etching, the contrast between these precipitates increases slightly, Figure 1-2i. In the false color image the following colors were chosen for the different phases: M23C6 – yellow, M(C,N) and MN, respectively – blue and cyan, respectively, G-phase – magenta, Figure 1-2j. In the outer metal sub-surface zone of the specimen in addition the carbide M6C (6.3%) was found which has, however, not been docu-mented photographically in the present case.

��00°F/�,�� h/�.� ksi (��°C/�� h/�0.0 MPa)After 7,217 h the number of creep pores had increased in the micro-

�� Atlas of Microstructures

structure, Figure 1-3a, which is also true for the secondary precipitates of M23C6, Figures 1-3a to 1-3j. However, the volume fraction of the primary and secondary M23C6 carbides remained constant with 7.8%. The fraction of the G-phase has slightly increased with 1.8%. One strik-ing feature is the bright, high Nb-containing dot-like inclusions in the G-phase which are residues of dissolved M(C,N), Figures 1-3c to 1-3h. The fraction of M(C,N) has decreased from 0.5 to 0.3% due to increased formation of G-phase. In the metal sub-surface zone M6(C,N) can replace the G-phase with a volume fraction of 12.5% and is most frequently directly in contact with the M23C6 carbide, Figures 1-3k to 1-3r. The composition of this phase can be taken from the table in Figure 1-3. In the DIC image this phase appears gray, while in the BSE image it can scarcely be distinguished from the matrix, Figures 1-3m and 1-3n. With increasing amounts of Ni this phase becomes brighter in the BSE image. In the ZnSe-coated condition the phase appears violet, Figures 1-3o and 1-3p, and in the false color image it had been assigned a red color, Figure 1-3r. The phase is resistant to V2A etchant, Figure 1-3q. The amount of M23C6 is slightly lower in the sub-surface area with 6.6% compared to the core area with 7.8%. In the outermost zone due to uptake of nitrogen the phase M2(C,N) appears which is slightly elevated in the metallographic section in the DIC image due to its higher hard-ness (not documented here).

1900°F/1038°C

��°F/3�� h/3.4 ksi (�050°C/3�� h/�3.4 MPa)Pores and cracks from the creep process can be clearly seen in the

microstructure, Figure 1-4a. In the core of the specimen the phases M23C6, G-phase and M6C could be identified by EPMA. In the light microscopic investigations and BSE image furthermore the phase M(C,N) could be found, Figures 1-4a to 1-4j. Since the G-phase and M6C could not easily be distinguished from each other in the contrasted condition both phases were taken together in the image analysis proce-dure. The fractions given for these two phases in the table of Figure 1-4

are based on an estimation. The primary and secondary M23C6 particles amount to 6.9%, the G-phase to about 0.7%, the M6C carbides to 1-2% and the M(C,N) particles to 0.4%. Due to the high amount of Nb the M6C carbides in the BSE image appear with a similar level of gray as the G-phase, Figures 1-4e and 1-4f, while M(C,N) has the bright-est appearance. In the sub-surface zone the amount of M23C6 increases slightly (7.4%), that of the M6C carbides reveals a significant increase (4.2%), while the G-phase no longer exists in these positions, Figures 1-4k to 1-4r. Again, directly under the outer part of the metal sub-surface zone M2(C,N) can be detected.

��00°F/�,0�3 h/�.0� ksi (�03�°C/�0�3 h/�4.4 MPa)With increasing exposure time the amount of pores and cracks from

creep has significantly increased, Figure 1-5a. The fraction of primary and secondary M23C6 (8.3%) and of the G-phase (3.0%) increase in the core of the specimen, while almost no M6C (< 0.03%) and M(C,N) (< 0.05%) can be found, Figures 1-5a to 1-5j. In the sub-surface zone the amount of M23C6 (8.8%) and of M6C (13.1%) has significantly increased. G-phase and M(C,N) are scarcely found there with a fraction of less than 0.02%, Figures 1-5k to 1-5r. Again, M(C,N) can only be found in the outermost zone of the sub-surface area.

��00°F/�,33� h/�.� ksi (�03�°C/�33� h/��.0 MPa)After exposure of more than 6,000 h a significant coarsening and

increase of the volume fraction of M23C6 (10.5%) and of M6C (7.5%) had taken place in the core and in the sub-surface area. The size of the pores has increased, Figures 1-6a and 1-6b. The secondary carbides have been dissolved or have become larger and blocky, Figures 1-6a to 1-6r. In the outermost zone of the sub-surface again M2(C,N) is detected which is slightly elevated in the LM-DIC image and appears somewhat darker than M23C6 and brighter than M6C, Figure 1-6d. In the SEM-BSE image this phase appears very dark so that it can be clearly distinguished

�3Atlas of Microstructures

from M6C, Figures 1-6u and 1-6v. After contrasting the phases by ZnSe M2(C,N) has a violet color similar to M6C with, however, a slightly brighter tone, Figure 1-6y. It was, however, not possible to distinguish between the two phases in the image analytical process due to an overlap in the color distributions in the red and blue range. In the false color image the M2(C,N) was colored green, Figure 1-6z.

��00°F/�0,�4� h/�.54 ksi (�03�°C/�0�4� h/�0.� MPa)After more than 10,000 h the amount of M23C6 has increased to

15% and that of M6C to 10.7%. The size of the carbides has increased even more, Figures 1-7a to 1-7j, and a larger number of M23C6 carbides is found than after 6,331 h. A large number of pores from creep are found which have agglomerated increasingly to cracks oriented perpendicular to the stress direction. In the sub-surface zone additionally M2(C,N) is found, Figures 1-7k to 1-7r.

2000°F/1093°C

�0��°F/335 h/�.0 ksi (��00°C/335 h/�3.� MPa)After 335 h of exposure at 2012°F (1100°C) the microstructure in

the core reveals about 7.6% M23C6 (primary and secondary carbides) and about 1.1% M(C,N), Figures 1-8h to 1-8j. Again a large number of creep pores and cracks are observed. In the sub-surface zone furthermore about 6.8% of M6C carbides are present which frequently contain resi-dues of the Nb-rich M(C,N) in their center, Figures 1-8k to 1-8r.

�000°F/�,�� h/�.� ksi (�0�3°C/�� h/��.4 MPa)With increasing exposure time under creep conditions the pores have

grown significantly, Figure 1- 9a. The M23C6 carbides have also grown in their size with the volume fraction in the core area of the specimen increasing up to 10.8%, Figures 1-9a and 1-9b. Besides a small frac-tion (0.3%) of M6C the microstructure also reveals the presence of the G-phase and of M(C,N). Due to an overlap in the colors the fractions could only be determined as a sum on the ZnSe-coated metallographic section, Figures 1-9g and 1-9h. It is estimated that about 1.2% are G-

phase, while 0.4% are M(C,N). The total fraction of these two phases was determined to be 1.6%. In the false color image the phases were partially re-drawn manually, Figure 1-9j. In the sub-surface zone again higher amounts of M6C (10.8%) are found beside M23C6 (10.3%), while the G-phase is missing, Figures 1-9k to 1-9r. The phase M2(C,N) is only found in the outermost area again, Figures 1-9s to 1-9z.

�0��°F/3,545 h/�.�5 ksi (��00°C/3545 h/�.� MPa)In the core area of the specimen the fraction of the significantly

coarsened M23C6 carbides has increased up to 12% and the fraction of M6C amounts to 3.8%, Figures 1-10a to 1-10j. Furthermore some very fine secondary precipitates of M23C6 are also found. The G-phase is no longer present. The fraction of M(C,N) amounts to 0.4%. In the sub-surface zone of the specimens about equal fractions of M23C6 and M6C are observed (8.8%). Furthermore the microstructure contains about 0.6% M(C,N) and M2(C,N), Figures 1-10k to 1-10r.

2050°F/1121°C

�05�°F/�� h/�.�0 ksi (��5°C/�� h/��.� MPa)After 277 h of exposure at the further increased temperature of

2057°F (1125°C) several rows of creep pores oriented perpendicular to the tensile direction are observed, Figure 1-11a. The microstructure in the core of the specimen contains 7.2% of carbides of the type M23C6 (primary and secondary carbides) and 1.1% of the type M(C,N), Figures 1-11c to 1-11j. Furthermore in the sub-surface zone, Figures 1-11k to 1-11r, 4.5% of M6C are observed while the fraction of M(C,N) has significantly decreased (0.2%). M2(C,N) is observed only locally in the outermost sub-surface area.

�050°F/�,�0� h/�.00 ksi (��°C/��0� h/�.� MPa)Compared to the shorter exposure time coarsening and an increase

in the number of M23C6 carbide to a fraction of 12.7% has taken place


in the core area of the specimen with a simultaneous decrease of the fine secondary carbides, Figures 1-12a to 1-12j. The formation of a signifi-cant fraction of carbides of the type M6C (9.2%) is associated with a dissolution of the carbides of the type M(C,N) whose fraction amounts only to 0.2%. In the metal sub-surface zone the carbide of the type M2(C,N) is found in a significant amount, Figures 1-12k to 1-12z.

��00°F/�4� h/�.50 ksi (��4�°C/�4� h/�0.3 MPa)After brief exposure at the highest temperature of 2100°F (1149°C)

a fraction of 8.6% of M23C6 primary and secondary carbides as well as of about 0.6% of M(C,N) is observed after 248 h with a small fraction of M6C (< 0.3%), Figures 1-13a to 1-13j. In the sub-surface zone the amount of M6C carbides is significantly higher (7.6%) with also the amount of M23C6 having increased (9.9%). Furthermore in the outer-most sub-surface zone which had been strongly oxidized and depleted in chromium M2(C,N) is also observed, Figures 1-13s to 1-13v.

Conclusions

The results of these investigations can be summarized as follows:The centrifugally cast material 35Cr/45Ni in the as-cast condition

consists of an austenitic matrix with eutectic areas of chromium-rich carbides of the type M7C3 and of Nb or Ti-rich particles of the type M(C,N) which have precipitated between the grains or along the grain boundaries respectively. Already after 200-300 h of exposure in the temperature range of 1800-2100°F (982 - 1149°C) a complete conver-sion of the carbides of the type M7C3 into chromium-rich carbides of the type M23C6 takes place with the latter type of carbides forming second-ary precipitates in addition. With increasing exposure time the number of these secondary carbides increases and also the fraction of the primary M23C6 carbides increases, accompanied by significant growth. At the same time additional M6C and/or Gphase are formed. Formation of the latter phases initiates at the phase boundary between M23C6 and matrix. M6C is rich in Cr, Ni and Nb, while the G-phase (Ni16Nb6Si7) contains

less Cr, but is rich in Ni, Si and Nb. Frequently the G-phase shows a seam of small round M23C6 secondary carbides. M6(C,N) shows a stronger tendency of formation in the metal sub-surface zone compared to the core area which may be explained by depletion processes due to outward diffusion of oxide-forming elements (e.g. Cr, Si) or by the uptake of nitrogen. The appearance of M6C seems to be stimulated at higher temperatures, i.e. it appears after shorter times. Furthermore the fraction increases with increasing exposure time. The G-phase is only found in the temperature range of 1800-2000°F (982 - 1149°C). In the lower part of this temperature range (1800°F (982°C)) this phase is still observed after 7,217 h, while at 2000°F (1093 °C) it has disappeared after 6,331 h. At 1900°F (1038°C) it can be found after 1,300 h, while after 3,545 h it no longer exists and seems to be replaced by M6C. At 2012°F (1100°C) this phase is not observed even after 335 h. In this case the lower Si content (1.52 mass-%) may have played a role9. An increased amount of silicon (> 1.6% mass-%) promotes the formation of G-phase by facilitating the conversion of NbC into G-phase. In the G-phase as well as in the M6C particles enrichment of Nb is observed which directs towards the dissolution of former MC10-15. The dissolution of MC with increasing exposure time is confirmed by a decrease of their volume fraction. In the cast state the fraction amounts to about 0.8%, while after short exposure times in which no G-phase has formed yet it amounts to 0.6-1.1%. When G-phase formation or after longer times G-phase and M6C formation start the fraction goes down to less than 0.02-0.4%. After several thousand hours above 1900°F (1038°C) the secondary carbides (M23C6) dissolve and form blocky, elongated parti-cles. Besides these coarse carbides of the type M23C6 where the particles sometimes also have grown together only coarse carbides of the type M6C and in the sub-surface zone of the type M2(C,N) are present. The G-phase is no longer found.


Table 10: Results of image analysis of Alloy HPMA Creep Rupture Stress SamplesSample_ ID Sample Temp. Temp °F Stress Stress Expo. area fraction in %

Time core outer zone / edge°C °F MPa Ksi hrs M(C,N) M7C3 M23C6 M6C M2(C,N) G-Phase M(C,N) M7C3 M23C6 M6C M2(C,N)

A234_018 HPMA-1 (c) as cast as cast 0.9 3.4 + ++A234_019 HPMA-2 927 1700 55.2 8.00 183 0.6 7.0 0.8 + +++ +A234_020 HPMA-3 927 1700 41.3 6.00 2,338 <0.3 8.3 0.7 + +++ +A234_021 HPMA-4 927 1700 36.6 5.30 6,072 0.8 8.0 0.8 ++ <0.3A234_022 HPMA-5 954 1750 30.7 4.45 8,359 0.8 8.7 + +++A234_023 HPMA-6 982 1800 41.2 5.97 177 0.6 7.0 0.4 5.3 4.9

4.7A234_024 HPMA-7 982 1800 24.2 3.51 2,436 1.2 7.9 0.4 6.4A234_025 HPMA-8 982 1800 26.9 3.90 6,478 0.5 8.1 0.5 0.5 6.1 1.0A234_026 HPMA-9 1000 1832 20.7 3.00 12,289 0.8 9.0 +A234_027 HPMA-10 1050 1922 24.0 3.48 137 0.7 9.5 1.3 0.4 5.6 9.8A234_028 HPMA-11 1066 1950 19.1 2.80 2,558 1.0 8.8 <0.2 0.9 7.1 5.1A234_029 HPMA-12 1050 1922 16.0 2.32 4,864 1.1 9.7 + +++A234_030 HPMA-13 1050 1922 14.3 2.08 11,778 0.8 13.1 1.7 0.8 0.0 15.2A234_031 HPMA-14 1100 2012 16.5 2.39 105 0.8 7.7 + +++ +A234_032 HPMA-15 1100 2012 12.4 1.80 2,714 0.8 9.6 + +++A234_033 HPMA-16 1100 2012 10.0 1.45 5,580 1.0 11.1 + ++++A234_034 HPMA-17 1100 2012 9.3 1.35 8,990 1.1 10.9 + ++++

Annotations: + few, ++ many, +++ very many, ++++ abundant


Alloy HPMAThe results of the image analytical determination of the volume

fractions of all specimens of this material are summarized in Table 10 (previous page).

As-CastFigures 2-1a to 2-1j show the as-cast condition after application of

the different imaging techniques. In the austenitic matrix two different types of precipitates can be distinguished which are located on the grain boundaries or in areas of the residual melt. As the EPMA analyses reveal these precipitates are Nb or Ti-rich carbides MC or carbon nitrides of the type M(C,N) and Cr-rich carbides of the type M7C3 which can be distinguished even in the DIC image due to their shape, color and size, Figures 2-1a to 2-1d. M(C,N) precipitates have different modifica-tions with varying contents of Nb, Ti and Cr. In the BSE image 2-1b the heavier elements Zr and Nb appear brighter than the lighter ones. Therefore the Ti-rich core of M(C,N) appears darker. A better contrast is achieved by ZnSe coating, Figures 2-1g and 2-1h, after which the M7C3 appear in purple, the Ti-rich cores of the M(C,N) particles in light blue to white and the highly Nb-containing species in blue. True-color imaging allows a reliable identification of the different phases and a quantification by the binary image, Figure 2-1j. The fractions of M7C3 amount to 3.4% and of M(C,N) to 0.9%.

Creep Rupture Samples

1700°F/927°C

��00°F/��3 h/�.00 ksi (��°C/��3 h/55.� MPa)Already after 183 h at a temperature of 1700°F (927°C) and a load

of 8.0 ksi the M7C3 has been converted into the more stable M23C6. Furthermore M23C6 secondary carbides are observed in the grain inte-rior, Figures 2-2b to 2-2d. The EPMA analyses reveal that the conversion of M7C3 into M23C6 is mainly due to an uptake of Ni from the matrix. The Ni content of M23C6 amounts to about 4-5 mass-%. As a further

phase containing Cr, Ni, Si and Nb this creep specimen exhibits the phase M6C which also contains some Fe. This phase appears as light gray in the DIC image and is partially encapsulated by M23C6, Figures 2-2c and 2-2d. Frequently in the M6C particles, which can hardly be distin-guished in the BSE image, Figures 2-2e and 2-2f, from the surrounding matrix bright Zrrich MC areas are also present. Coating of the specimen with ZnSe facilitates the detection of the M6C phase (which can easily be overlooked in the BSE image) by changing the color to violet. This is illustrated by the comparison of Figure 2-2f and Figure 2-2h. In the false color image, Figure 2-2j, the M23C6 (7.0%) are colored yellow, the M6C (0.8%) are colored red and the M(C,N) (0.6%) are colored blue.

��00°F/�,33� h/�.00 ksi (��°C/�33� h/4�.3 MPa)Even after more than 2,000 h of exposure the microstructure has

not changed significantly, Figures 2-3a to 2-3r. In the core about 8.3% M23C6, 0.7% M6C and less than 0.3% MC are observed. The decrease in the amount of MC is probably due to the fact that according to chemi-cal analysis this specimen contains less Nb and no Ti, as well as a lower amount of C, so that in the end fewer M(C,N) particles were formed. The sub-surface zone has a similar appearance of the microstructure.

��00°F/�,0�� h/5.30 ksi (��°C/�0�� h/3�.� MPa)Even after 6,000 h of exposure no significant changes in the micro-

structure can be found. The fraction of M23C6 in the core area amounts to 8.0%, while the amount of M(C,N) has increased to 0.8% due to the chemical composition of the specimen, Figures 2-4a to 2-4f. In the sub-surface zone of the specimen some M6C (< 0.3%) had precipitated, Figures 2-4k to 2-4r.

��50°F/�,35� h/4.45 ksi (�54°C/�35� h/30.� MPa)After more than 8,000 hours at a slightly increased temperature the

amount of M23C6 has increased slightly from 8.0% to 8.7% which is probably due to an increased precipitation of secondary carbides, Figures 2-5a to 2-5r. The amount of M(C,N) of 0.8% has remained constant


which is consistent with the identical compositions and phases formed in the specimens numbers HP-MA-4 and HP-MA-5.

1800°F/982°C

��00°F/�� h/5.�� ksi (��°C/�� h/4�.� MPa)In the austenitic matrix of the specimen again about 7.0% M23C6

(primary and secondary carbides) as well as 0.6% M (C,N) have been found, Figures 2-6a to 2-6r. The outermost sub-surface zone additionally contains about 5% M6C and M2(C,N) respectively (not documented here). The preferred formation of M6C in the sub-surface area is presum-ably due to the relatively high amount of Nb in the specimen and the outward diffusion of Cr and Si which are needed for the formation of the oxide scale. Compared to M23C6 the carbide M6C is less rich in Cr and Fe, but contains more Ni, Si and Nb. The formation of M2(C,N) which is formed in the outermost sub-surface zone is probably due to outward diffusion of Cr and nitrogen uptake from the environment. The matrix in the metal sub-surface zone was strongly depleted in Cr around these M2(C,N) particles.

��00°F/�,43� h/3.5� ksi (��°C/�43� h/�4.� MPa)After 2,436 h the amount of secondary carbides evidently increases

slightly which is reflected by an increase of M23C6 to 7.9%, Figures 2-7a to 2-7j. The high amount of M(C,N) of 1.2% is due to the composi-tion of the material, cf. Table 4. Again in the metal sub-surface zone the formation of M6C and M2(C,N) is observed, Figures 2-7k to 2-7z.

��00°F/�,4�� h/3.�0 ksi (��°C/�4�� h/��.� MPa)After more than 6,000 h besides M23C6 (8.1%) and M(C,N) (0.5%)

smaller amounts of M6C (0.5%) are also found in the core area, Figures 2-8a to 2-8j. In the metal sub-surface zone the amount of M6C is twice as high as in the core, Figures 2-8k to 2-8r.

��3�°F/��,�� h/3.0 ksi (�000°C/�� h/�0.� MPa)

After more than 12,000 h a coarsening of the primary carbides and a coagulation of secondary carbides have taken place, Figures 2-9a to 2-9r. The fraction of M23C6 amounts to 9.0%, M6C is observed neither in the core nor in the metal subsurface zone. The fraction of M(C,N) amounts to 0.8% and is clearly higher than in the case of specimen HP-MA-8 where in the core M6C had also been found.

1920°F/1050°C

��°F/�3� h/3.4� ksi (�050°C/�3� h/�4.0 MPa)At this further elevated temperature a significant increase in the

number of creep pores and cracks can be observed, Figure 2-10a. The M23C6 primary carbides seem to have coarsened and the secondary carbides show a needle-like and blocky appearance, Figures 2-10c to 2-10r. Their fraction amounts to 9.5% in the core area with that of M(C,N) amounting to 0.7%. Furthermore the core area contains 1.3% M6C. The metal sub-surface zone shows the following fractions: M23C6 – 5.6%, M23C6 – 0.4% and M6C and M2(C,N) (as a sum) – 9.8%, Figures 2-10k to 2-10r.

��50°F/�,55� h/�.�0 ksi (�0��°C/�55� h/��.� MPa)After more than 2,500 h the amount of creep damage has further

increased, Figure 2-11a. In the core area the beginning of M6C forma-tion (< 0.2%) has been found. The fraction of M23C6 amounts to about 8.8% and that of M(C,N) to about 1.0%, Figures 1-11a to 1-11j. The microstructure of the sub-surface zone contains the following phases: M(C,N) – 0.9%, M23C6 – 7%, and M6C and M2(C,N) as a sum – 5%, Figures 2-11a to 2-11r.

��°F/4,��4 h/�.3� ksi (�050°C/4��4 h/��.0 MPa)After almost 5,000 h of exposure the amount of M23C6 carbides has

increased up to 9.7%. The fraction of M(C,N) amounts to about 1%. M6C and M2(C,N) are neither observed in the core nor in the metal sub-surface zone, Figures 2-12a to 1-12r.


��°F/��,�� h/�.0� ksi (�050°C/�� h/�4.3 MPa)After almost 12,000 h a significant coarsening and an increase in

the amount of M23C6 carbides has taken place in the core area of the specimen with about 13.1%. M2(C,N) has spread from the metal sub-surface zone into the core area, Figures 2-13c to 2-13j. Due to its higher hardness M2(C,N) can easily be distinguished from the other carbides in the DIC image. Even in the BSE image it shows a good contrast, Figures 2-13e and 2-13f. EPMA investigations reveal that this carbonitride is almost free of Ni and has the composition (Cr, Fe)2 (C,N), cf. table in Figures 2-13a and 2-13b. Values for the ratio between carbon and nitrogen cannot be given here since C and N had not been analyzed. Although the particles of M2(C,N) are rather coarse, which should facil-itate their identification, it is not easy to distinguish them from M6C in the ZnSe-coated section, Figures 2-13e to 2-13r. The false color image, Figure 2-13j, was constructed by image manipulation in which the color green was assigned to M2(C,N). Besides about 0.8% M(C,N) about 1.7% M6C plus M2(C,N) were measured in the core area, Figures 2-13a to 2-13j. In the metal sub-surface zone no M23C6 was found. Besides about 15% M2(C,N) about 0.8% M(C,N) were present, Figures 2-13k to 2-13r.

2012°F/1100°C

�0��°F/�05 h/�.3� ksi (��00°C/�05 h/��.5 MPa)Compared to specimen HP-MA-10 (1922°F/137 h) (xxx) the

amount of M23C6 carbides is lower with the precipitates being coarser, Figure 2-14a to 2-14r. Almost no differences between the core area and the metal sub-surface zone are observed with 7.7% M23C6 beside 0.8% M(C,N).

�0��°F/�,��4 h/�.�0 ksi (��00°C/��4 h/��.4 MPa)After 2,714 h the fraction of M23C6 carbides has increased due to an

increase of secondary precipitates up to 9.6%. The fraction of M(C,N) of 0.8% remained about constant. The microstructure in the core and in the metal sub-surface zone is identical, Figures 2-15a to 2-15r.

�0��°F/5,5�0 h/�.45 ksi (��00°C/55�0 h/�0.0 MPa)With increasing exposure time the amount of M23C6 increases further

(11.1%). The number of the secondary carbides decreases since the very fine particles combine with the larger ones to form coarse particles and some of them seem to agglomerate with the primary carbides, Figures 2-16a to 2-16r. The fraction of M(C,N) lies around 1%. Occasionally high Nb-containing needles are found which most probably are nitrides, Figure 2-16b. The thickness of these needles is less than 1 μm and their length reaches up to 100 μm. In the DIC image they appear gray and after ZnSe coating violet.

�0��°F/�,��0 h/�.35 ksi (��00°C/��0 h/�.3 MPa)After almost 9,000 h of exposure time no visible changes in the

microstructure are observed compared to the former specimen, Figure 2-17a to 2-17r. Even the fractions of M23C6 (10.9%) and of M(C,N) with 1.1% are comparable. Again highly Nb-containing precipitates are found, Figures 2-17b and 2-17s.

Conclusions

The results of the investigations can be summarized as follows:The centrifugally cast material HP-MA (microalloy) in its cast state

consists of an austenitic matrix with a eutectic structure at interdendritic sites or along grain boundaries respectively, formed by Crrich carbides of the type M7C3 and Nb or Ti-rich carbonitrides or carbides of the type M(C,N). Already after 100-200 h of exposure in the temperature range of 1700-2012°F (927 - 1100°C) a complete conversion of the carbides of the type M7C3 into Cr-rich carbides of the type M23C6 takes place. The M23C6 carbides form secondary precipitates. With increas-ing exposure time the number of secondary carbides increases and later on also the number of primary M23C6 carbides rises accompanied by a strong growth in size. The fraction increases from about 7% to a maxi-mum of 11-13%. At the same time some M6C is precipitated in the core


area, while this is more pronounced in the metal sub-surface zone. At 1700°F (927°C) the M6C precipitates with an amount of less than 1% are dissolved after times longer than 2,500 h in the core of the speci-men and exist in the metal sub-surface area up to about 6,000 h where, however, only an amount of less than 0.3% can still be detected. After more than 8,000 h no more M6C can be found. At the higher tempera-ture of 1800°F (982°C) M6C can be found only in the metal sub-surface zone at times up to 2,500 h. After 6,500 h traces (< 0.5%) of M6C are observed in the core area, while the fraction in the metal sub-surface zone decreases significantly. Due to the uptake of nitrogen M2(C,N) is also formed whose fraction was measured as part of the sum with M6C. After 12,000 h at 1832°F M6C is neither found in the core nor in the metal subsurface zone. At temperatures above 1900°F (1038°C) small amounts of M6C (< 0.2-1.3%) are observed in the core after exposure times of 2,500 h which becomes dissolved after longer times. The forma-tion of M6C and M2(C,N) occurs only in the metal sub-surface zone to a significant extent with their fraction decreasing markedly after 2,500 h due to outward diffusion processes. However, after almost 5,000 h M6C and M2(C,N) are observed neither in the core nor in the metal sub-surface zone. After almost 12,000 h the extension of M2(C,N) is marked and reaches from the sub-surface zone into the core of the specimen. In the core only single precipitates of M6C are observed besides M23C6 and M(C,N), while in the outer part of the specimen only M2(C,N) and M(C,N) are present. At the maximum temperature of 2012°F (1100°C) the increase of the fraction of M23C6 with time is very marked. M6C and M2(C,N) are not present in the core and are observed only after short exposure time (100 h) and sporadically in the metal sub-surface zone. After longer exposure times in the core area some highly Nb-contain-ing needle-shaped phases are formed. In the metal sub-surface zone an increasing depletion of oxide-forming elements takes place so that finally M6C and M2(C,N) are dissolved completely. In none of the specimens investigated was the G-phase detected.

Alloy HP ModifiedThe results of the image analytical determination of the volume frac-

tions of the precipitates are summarized in Table 11.

As-CastFigures 3-1a to 3-1j show the as-cast condition after application of

the different imaging techniques. The microstructure is comparable to that of the as-cast condition of the material HPMA. The precipitates along the grain boundaries or in the interdendritic residual melt areas are the Cr-rich carbides of the type M7C3 (3.9%), which appear bright gray in the DIC image, and Nb-rich carbides of the type M(C,N) with 0.9% which appear slightly red in the DIC image, Figures 3-1c and 3-1d. In the ZnSe-coated state M7C3 appears purple and M(C,N) blue, Figures 3-1g and 3-1h.

Creep Rupture Samples

1675°F/913°C

��5°F/�5� h/�.50 ksi (��3°C/�5� h/44.� MPa)At the lowest temperature of 1675°F (913°C) after 659 h in the

core primary carbides of the type M23C6 are found together with elon-gated secondary carbides of the same type (together 8%). Furthermore some carbides of the type M6C (0.5%) and its secondary precipitates are observed together with about 1.8% G-phase and 0.3% M(C,N), Figures 3-2a to 3-2j. The M(C,N) precipitates besides a high amount of Nb contain also Zr, see table in Figures 3-2a and 3-2b. Presumably the formation of G-phase was favored by the high amount of Si of 1.67 mass-%, see Table 6. In the DIC image it was not easy to distinguish between the G-phase and the carbides of the type M6C since both of them appear gray, Figures 3-2c and 3-2d. In the BSE image the G-phase appears bright due to its high Nb content, while M6C does not differ much from the matrix, Figures 3-2e and 3-2f. The M(C,N) particles appear bright in the BSE image. These particles can, however, be distin-


Table 11: Results of image analysis of Alloy HP Modified Creep Rupture Stress SamplesSample_ID Sample Temp. Temp °F Stress Stress Expo. area fraction in %

Time core outer zone / edge°C °F MPa Ksi hrs M(C,N) M�C3 M�3C� M�C M�(C,N) G-Phase M(C,N) M�C3 M�3C� M�C M�(C,N) G-Phase

A234_036 HPNb-1 (b) as cast as cast 0.9 3.9 + ++A234_038 HPNb-2 913 1675 44.6 6.5 659 0.3 8.0 0.5 1.8 (+) +++ + +A234_039 HPNb-3 927 1700 58.8 8.5 59 1.0 8.0 0.7 + +++ +A234_040 HPNb-4 955 1750 35.2 5.1 794 0.4 7.5 2.4 (+) (+) +++ ++A234_041 HPNb-5 970 1778 35.5 5.2 286 0.5 6.7 1.1 + +++ ++A234_042 HPNb-6 970 1778 26.0 3.8 1,185 0.4 7.1 0.7 (+) +++ +A234_043 HPNb-7 982 1800 37.4 5.42 191 3.0 7.6 ++ +++ (+) (+)A234_044 HPNb-8 982 1800 37.4 5.42 4,467 2.2 6.4 ++ +++ (+)A234_045 HPNb-9 982 1800 37.4 5.42 7,833 1.7 7.1 ++ +++ (+) (+)A234_046 HPNb-10 982 1800 19.3 2.8 10,637 0.4 11.1 3.9 (+) ++++ ++ +A234_047 HPNb-11 1010 1850 24.8 3.6 707 0.7 7.8 0.9 + +++ ++ +A234_048 HPNb-12 1038 1900 15.8 2.3 2,555 0.6 8.4 2.5 + +++ ++ +A234_049 HPNb-13 1038 1900 13.1 1.9 5,373 0.6 9.5 1.0 + ++++ ++ +

Annotations: (+) very few, + few, ++ many, +++ very many, ++++ abundant


guished from the G-phase due to their blocky shape. After ZnSe coating the M(C,N) particles appear in light blue to white, the G-phase in blue, the M6C particles in violet and the M23C6 carbides in magenta and are thus easy to distinguish. The G-phase particles and the M6C particles had to be measured together in the image analytical investigations since it was not easy to distinguish between the two of them and to measure their fractions separately. The fractions of the two phases were estimated manually. The false color image in Figure 3-2j needed some manual adjustment and the phases were colored as follows: M23C6 – yellow, G-phase – magenta, M(C,N) – blue and cyan, M6C – red. Etching with V2A etchant improves the visibility of the precipitates in the matrix, Figure 3-2i. The microstructure in the sub-surface zone is not different from that in the core, Figures 3-2o to 3-2r.

1700-1750°F/913-955°C

��00°F/5� h/�.5 ksi (��°C/5� h/5�.� MPa)After this short exposure time at 1700°F (927°C) the microstructure

is full of secondary precipitates. The M23C6 carbides (8.0%) show typi-cal eutectic appearance, Figures 3-3a to 3-3r. Besides a small amount of M6C (0.7%), some M(C,N) are also observed with about 1% in the core area as well as in the metal sub-surface zone.

��50°F/��4 h/5.�0 ksi (�55°C/��4 h/35.� MPa)After this increased exposure time at 1750°F (954°C) the amount of

M6C in the core area has slightly increased compared to specimen HP-Nb-3, while the amount of M(C,N) has clearly decreased, Figures 3-4a to 3-4j. Possibly some G-phase is present in the M6C carbides which, however, could not be evidenced by EPMA. The bluish coloring in the violet M6C seemed to indicate their presence, Figure 3-4a. Since the chemical composition of this specimen is identical to specimen HP-Nb-2, the appearance of the G-phase would have been expected here. In the image analysis this phase was measured together with M6C or M(C,N) so that their fractions may appear slightly too high. In the sub-surface zone

of the specimen residues of dissolving M(C,N) are frequently detectible due to their blocky shape and are embedded into M6C, Figures 3-4k to 3-4r. Bluish colorings which would indicate the presence of G-phase were not detected here.

1778°F/970°C

��°F/�� h/5.�5 ksi (��0°C/�� h/35.5 MPa)After this relatively short exposure time at 1778°F (970°C) the core

area contains 6.7% M23C6, 1.1% M6C and 0.5% M(C,N), Figures 3-5a to 3-5j. Despite the relatively high amount of Si of 1.79 mass-% the G-phase could not be identified reliably. The metal sub-surface zone does not show any differences compared to the core area, Figures 3-5k to 3-5r.

��°F/�,��5 h/3.�� ksi (��0°C/��5 h/��.0 MPa)After this extended exposure time of almost 1,200 h no differences

with regard to the volume fractions of the phases identified by EPMA are observed compared to specimen HP-Nb-5, Figures 3-6a to 3-6r. The G-phase is not detected despite the high Si amount of 1.79 mass-%.

1800°F/982°C

��00°F/�� h/5.4� ksi (��°C/�� h/3�.4 MPa)The specimens HP-Nb-7 to HP-Nb-9 possess a high amount of Nb

(1.2 mass-%) which has a significant effect on the microstructure. Besides M23C6 a large number of M(C,N) are precipitated in very fine structures and a eutectic appearance, Figures 3-7a to 3-7r. Their volume fraction amounts to 3% with an amount of 7.6% of M23C6. In the outermost zone of the specimen some M6C and M2N are encountered.

��00°F/4,4�� h/5.4� ksi (��°C/44�� h/3�.4 MPa)After an exposure time of 4,467 h the microstructure has not changed

compared to specimen HPNb- 7, Figures 3-8a to 3-8r. The volume frac-tions of the carbides determined in the image analytical investigations


are about equal to the former specimen.

��00°F/�,�33 h/5.4� ksi (��°C/��33 h/��.3 MPa)Even after almost 8,000 h no changes in the microstructure can be

found compared to the short exposure times, Figures 3-9a to 3-9r. Again the volume fractions of the carbides are unchanged.

��00°F/�0,�3� h/�.�0 ksi (��°C/�0�3� h/��.3 MPa)Compared to specimens HP-Nb-7 to HP-Nb-9 this specimen

HP-Nb-10 contains only 0.74 mass-% Nb, which is reflected by the appearance of the M(C,N) precipitates. The microstructure does not show the typical eutectic characteristics, Figures 3-10a to 3-10r. The long exposure time of more than 10,000 h has led to a coarsening of the primary carbides of M23C6. The fraction of M(C,N) plus M23C6 amounts to 11%. In addition about 4% M6C and 0.4% M(C,N) were formed. In the metal sub-surface zone the amount of M6C is significantly higher and due to the uptake of nitrogen some M2(C,N) is observed, Figures 3-10k to 3-10r.

1850-1900°F/1010-1038°C

��50°F/�0� h/3.�0 ksi (�0�0°C/�0� h/�4.� MPa)This specimen contains a high amount of Si with 1.7 mass-%. The

G-phase, however, could not be detected here. The following precipi-tates were observed: M23C6 – 7.8%, M2(C,N) – 0.7% and M6C – 0.9%. M(C,N) or their undissolved residues are frequently observed in the core of M6C particles, Figures 3-11a to 3-11j. In the outermost sub-surface zone additionally some M2(C,N) is present, Figures 3-11k to 3-11r.

��00°F/�,555 h/�.30 ksi (�03�°C/�555 h/�5.� MPa)After 2,555 h at slightly increased temperature growth of the M23C6

carbides had taken place. The number of the secondary carbides has decreased. The volume fraction of M23C6 has increased to 8.4% and that of M6C to 2.5%, while the fraction of M(C,N) has remained almost constant with 0.6%, Figures 3-12a to 3-12r. In the outermost zone some M2(C,N) is also present, which is colored green in the false color

image in Figure 3-12r.

��00°F/5,3�3 h/�.�0 ksi (�03�°C/53�3 h/�3.� MPa)After doubling the exposure time at a test temperature of 1900°F

(1038°C) the M6C carbides have become partially dissolved, while again increasingly M23C6 secondary carbides were precipitated, Figures 3-13a to 3-13r. Their volume fraction of M23C6 increases up to 9.5%, while that of M6C decreases to 1%. The amount of M(C,N) with 0.6% remains stable.

Conclusions

The results of these investigations can be summarized as follows:The centrifugally cast material HP-Nb (modified alloy) in its cast

version consists of an austenitic matrix with eutectic structures of the Cr-rich carbides of the type M7C3 and of Nb and Zr carbides of the type M(C,N). These structures are found along grain boundaries and in interdendritic areas. Already after an exposure time of 59-700 h and a temperature range of 1700-1900°F ( 927 -1038°C) a complete conver-sion of the carbides of the type M7C3 into Cr-rich carbides of the type M23C6 takes place with the latter also forming secondary precipitates. In addition in specimens with Nb contents between 0.74 and 0.83 mass-% M6C carbides are precipitated. The amount of M6C seems to increase in the lower temperature range up to 1750°F (954°C) with increasing exposure time. At the higher temperatures above 1778°F (970°C) rather a dissolution of M6C seems to take place with increasing exposure time. An exception is found for the specimen which had been exposed for 10,637 h at 1800°F (982°C) which had shown the highest amount of M6C (3.9%). Higher amounts of Nb (1.2 mass-%) seem to promote the precipitation of finely dispersed NbC as eutectic structures evidently stabilizing the microstructure and suppressing the formation of M6C. After longer exposure times at temperatures of 1800°F (982°C) the amount of M23C6 increases after times higher than 8,000 h and


at 1900°F (1038°C) after times higher than 2,500 h from about 8% to maximum 11%. The G-phase was observed only for one specimen of this material which was subjected to the lowest test temperature of 1675°F (913°C) for 659 h. At higher temperatures for specimens with comparable or even higher Si content the G-phase could not be detected with confidence.


Table 7: Exposure Conditions of Service-Exposed SamplesTube Skin Tube Skin Hoop Hoop Expo.

Sample_ID Sample ID Temp. Temp. Stress Stress Time DonatedDechema MTI Alloy °F °C ksi MPa hrs byA234_050 SE-A HPNb 1630 888 2.12 14.60 108,306 MetalTekA234_051 SE-B HPNb 1607 875 2.12 14.60 108,306 MetalTekA234_052 SE-C HPNb 1650 899 0.72 4.96 75,000 EastmanA234_053 SE-D HPNb ~1650 899 1.16 8.00 148,920 EastmanA234_054 SE-E HPNb ~2350 1288 1.30 9.96 35,000 APCIA234_055 SE-F HPMA ~1800 982 1.79 12.30 65,664 APCIA234_056 SE-G HPMA 1673 912 2.04 14.00 94,900 APCIA234_057 SE-H HPNb 1661 905 1.98 13.60 131,000 SyncrudeA234_058 SE-I HPNb 1661 905 1.98 13.60 143,000 SyncrudeA234_059 SE-J HPNb 1661 905 1.98 13.60 143,000 Syncrude

Table 8: Chemistry of Service Exposed SamplesSample_ID Sample ID Composition (wt%)Dechema MTI Cr Ni Fe Mn Si C Nb N OtherA234_050 SE-A 24.5 34.1 bal 0.40 0.90 0.45 0.80 0.047 Ti,Zr & W<0.1A234_051 SE-B 24.5 34.0 bal 0.34 0.90 0.39 0.75 0.060 Ti,Zr & W<0.1A234_052 SE-C 24.9 31.7 bal 0.68 1.10 0.42 0.74 0.074 Ti,Zr & W<0.1A234_053 SE-D 24.5 34.5 bal 0.83 1.90 0.40 1.40 0.062 Ti,Zr & W<0.1A234_054 SE-E 25.3 32.6 bal 1.20 2.50 0.40 1.20 0.074 Ti,Zr & W<0.1A234_055 SE-F 24.9 32.8 bal 1.00 1.30 0.42 0.98 0.038 Ti & W>0.1, Zr<0.1A234_056 SE-G 24.7 34.0 bal 0.70 0.90 0.57 0.50 0.104 Ti,Zr & W<0.1A234_057 SE-H 23.9 31.8 bal 0.50 1.10 0.42 1.20 0.128 Ti,Zr & W<0.1A234_058 SE-I 24.4 33.3 bal 0.60 1.40 0.42 1.20 0.090 Ti,Zr & W<0.1A234_059 SE-J 24.3 33.4 bal 0.57 1.40 0.39 1.20 0.102 Ti,Zr & W<0.1

Service-Exposed SamplesBelow the history of the Service-Exposed Tubes is described in detail. Exposure conditions and chemical composition are summarized

in Tables 7-8.


Sample SE-D was extracted from the failed tube at the same eleva-tion as the through-wall rupture. However, the sample was removed 180 degrees around the circumference from the rupture in an effort to obtain a microstructure representative of long-term ageing.

Sample SE-E:The HP modified reformer tube was in service for approximately

four years prior to rupture (see Table 8). The ruptured area of the tube had a maximum diametral growth of only 2.3%. Metallographic analysis revealed numerous, tight interdendritic cracks running nearly through wall with minimal creep voids. Higher magnification examina-tion revealed a lamellar-like phase adjacent to the interdendritic cracks. The tight intergranular cracks and lamellar phase are characteristic of incipient melting. Incipient melting is liquation of low melting point constituents along the dendrite boundaries. Incipient melting of cast HP alloys generally occurs at approximately 2350°F (1288°C) or 100°F (56°C) below the bulk melting temperature of approximately 2450°F (1343°C).

Samples SE-A & SE-B:These samples were removed from the same furnace after approxi-

mately 108,306 hours of service. The available history indicated that the furnace had 53 cold start-ups and 70 hot restarts over the service life. Tube samples SE-A & SE-B were removed from different tube segments and elevations within the furnace resulting in a slight difference in esti-mated tube skin temperatures (see Table 7).

Post service analysis found that samples SE-A and SE-B both had maximum diametral growths of less than 1%. Ambient temperature tensile testing of the aged material found the following:

SE-A SE-BUTS 66.9 Ksi / 462 MPa 65.8 Ksi / 454 MPa

Yield Strength 37.6 Ksi / 259 MPa 40.1 Ksi / 277 MPaElongation 5.9% 3.0%

Sample SE-C:This sample was removed within 6 inches (15.2 cm) from the outlet

end of an HP-modified tube. The tube had been in service approxi-mately 75,000 hours with no operational anomalies. Sample was included to observe the microstructural changes at relatively mild oper-ating conditions.

Sample SE-D:This tube failed approximately four months after a scheduled plant

outage. The failure was likely the result of debris inadvertently left in the tube during the outage. The debris caused channeling of the feed gas resulting in a local hot spot and subsequent rupture. Immediately adja-cent to the rupture, the secondary carbides were completely dissolved with notable coarsening of the primary carbides. The microstructural features suggested that the metal temperature at the ruptured area of the tube exceeded 2012°F (1100°C).


Micrograph of fine, interdendritic cracks and lamellar-like phase.(oxalic acid electrolytic, 500X)

The plant had a history of operational problems that tended to crush the reformer tube catalyst and promote channeling of the feed gas. The tube rupture resulted from a short-term temperature excursion. This sample was included to show the microstructural features as a result of incipient melting.

Sample SE-F:During an outage after 7.5 years of service, diametral growths

in excess of 10% were measured at an isolated area in the furnace. Metallographic analysis revealed only aligned creep voids with no creep fissures despite the large diametral growth. During normal operation, the tube metal temperature was in the 1650-1690°F (899-921°C) range. The microstructural features (i.e. carbide coalescence) and review of historical operating data indicated that the tube metal temperatures likely exceeded 1800°F (982°C) at an isolated area of the furnace for approximately 6 months leading up to the outage.

Chemical analysis found deficiencies in micro-alloying and/or carbon concentrations in many of the tube segments that experienced large diametral growths. Remaining life assessment via the Omega creep strain methodology was performed finding the creep strain rates as least 60 times greater than the average for HP micro-alloyed material in the Materials Property Council (MPC) database. The alloying deficiencies resulted in the material being highly strain tolerant, likely capable of 15% diametral growth prior to rupture, but with poor creep strength. The remaining life testing indicated that the material had essentially no remaining life at the normal operating conditions.

Sample SE-G:The HP micro-alloyed tube was removed from the furnace for

destructive remaining life assessment. The tube had been in service approximately 95,000 hours with a design life of 100,000 hours (typical operating conditions listed in Table 8). The tube had a maximum diam-etral growth of 5%. Metallographic analysis found the onset of creep void alignment but no creep fissure formation suggesting considerable remaining life.

Remaining life assessment was performed via the Omega creep strain method. Despite the actual service life reaching the original design life, the measured strain rates and calculated Omega indicated the tubes have an estimated remaining life of nearly 17 years at the normal operating conditions. The estimated tube life decreased approximately 44% with each 25 °F (14°C) increase in tube metal temperature.

Sample SE-H:This HP-modified sample was extracted from a tube removed from a

furnace after approximately 131,000 hours of operation (typical operat-ing conditions in Table 8). The sample was removed from the fired side of the tube approximately 16 inches (40.6 cm) above the floor of the top-fired, downflow reformer.

Eddy current inspection of the tube immediately prior to removal indicated that the tube had a 30% through-wall indication and 2.0-


2.4% diametral growth. Radiography could not confirm the indication. Subsequent metallography revealed small, linear indications, especially at the tube ID, where creep voids began to coalesce and align. Creep voids were present throughout the cross section but with little to no creep void alignment. Remaining life based on metallography was likely considerable, perhaps another 50,000 hours at the current design condi-tions. Caliper and strapping measurements found up to 5.6% diametral growth over the nominal diameter.

Upon removal from service, tube remaining life assessment was performed using the Omega creep strain methodology. Even though the tube had surpassed the design life by nearly 50%, the calculated Omega value and measured creep strain rate values indicated that the tube still had 20,000 to 60,000 hours of remaining life at the current operating conditions. The obtainment of life well in excess of the original design life was likely due to the operating temperature being approximately 45°F (25°C) lower than the 1710°F (932°C) design temperature.

Samples SE-I & SE-J:Samples SE-I & SE-J were obtained from a tube removed from the

same reformer as sample SE-H (above). However, samples SE-I & SE-J had increased service time of approximately 143,000 hours. Sample SE-I was extracted approximately 24 inches (61 cm) above the reformer floor while sample SE-J was extracted from the same tube but 84 inches (213.4 cm) above the reformer floor. The samples were included in an effort to observe any microstructural differences resulting from the slight eleva-tion changes within the furnace. Strapping and caliper measurements indicated that sample SE-I had a diametral growth of 2% while sample SE-H had 3.5% diametral growth. Tube metal temperature measure-ments taken at these two elevations indicated a difference of 9 °F (5°C), consistent with the diameter change. Omega creep strain remaining life assessment has not yet been completed in these samples.


Service-Exposed SamplesThe results of the image analytical determination of the volume frac-

tion of the precipitates for the service-exposed tubes of the centrifugally cast materials are summarized in Table 12.

Table 12: Results of image analysis of Service Exposed SamplesSample_ID Sample Material Tube Skin Temp. Hoop Stress Expo. area fraction in %

Time core outer zone / edge°F °C Ksi MPa hrs M(C,N) M23C6 M6C M2(C,N) G-Phase M(C,N) M23C6 M6C + M2(C,N) G-Phase

A234_056 SE-G HPMA 1673 912 2.04 14.0 94,900 0.3 12.9 1.6 + + +++ ++ +A234_055 SE-F HPMA 1800 982 1.79 14.0 65,664 <0.2 6.2 2.0 + +++ +A234_051 SE-B HPNb 1607 875 2.12 12.30 108,306 7.0 1.5 1.2 +++ + +A234_050 SE-A HPNb 1630 888 2.12 14.60 108,306 0.6 9.2 0.8 0.4 (+) +++ ++ (+)A234_052 SE-C HPNb 1650 899 0.72 4.96 75,000 8.0 2.2 1.0 +++ ++ ++A234_053 SE-D HPNb 1650 899 1.16 8.00 148,920 1.6 5.5 0.2 4.6 6.2A234_057 SE-H HPNb 1661 905 1.98 13.60 131,000 0.1 7.4 2.0 0.5 +++ ++ +A234_058 SE-I HPNb 1661 905 1.98 13.60 143,000 6.5 2.2 1.5 7.9 6.2A234_059 SE-J HPNb 1661 905 1.98 13.60 143,000 7.9 2.5 1.6 7.8 6.5A234_054 SE-E HPNb 2350 1288 1.30 9.96 35,000 1.3 4.0 + + (+)

Annotations: (+) very few, + few, ++ many, +++ very many


Alloy HPMA

��3°F/�4,�00 h/�.04 ksi (��°C/�4�00 h/�40 MPa)This specimen contains a very large number of M23C6 primary and

secondary carbides, Figures 2- SE-Ga to 2-SE-Gr. Their volume fraction amounts to 12.9%. Furthermore in the center of the tube wall as well as in the metal sub-surface zone carbides and nitrides of the type M6C and M2(C,N) are present. The sum of these carbides amounts to about 1.6%. The M(C,N) have evidently been almost completely dissolved or were oxidized and appear white in the cross-section which had been coated with ZnSe. Their fraction amounts to 9.3%. The G-phase could not be identified in this specimen. Besides the precipitates a large number of creep pores could be observed. In the BSE image the M2(C,N) appear bright due to the high Nb content, see Table in Figures 2-SE-Ga and 2-SE-Gb. The high amount of N (0.1 mass-%) could explain the forma-tion of M2N, which contains nearly 50 atomic-% Nb. As Nb is bound to N it is not available for the formation of G-Phase. High N-contents seem to prevent G-Phase formation.

��00°F/�5,��4 h/�.�� ksi (��°C/�5��4 h/��.3 MPa)After operation for 7.5 years at 1800°F (982°C) the microstructure

of the tube contains a high level of creep pores and shows M23C6 primary and secondary carbides all over the entire cross-section with a fraction of about 6.2%. In addition less than 2.0% G-phase due to the high Si content of 1.4 mass-% and the low N content of 0.04 mass-% as well as residues of M(C,N) with about less than 0.2% are observed, Figures 2-SE-Fa to 2-SE-Fr.

Alloy HP Modified

��0�°F/�0�,30� h/�.�� ksi (��5°C/�0�30� h/�4.� MPa)After an operation time of more than twelve years at 1607°F (875°C)

only a few creep pores are observed. The carbides are arranged in a distinctive eutectic structure, Figure 3-SE-Ba to 3-SE-Br. Besides about 7.0% M23C6 primary and secondary carbides about 1.2% G-phase and


1.5% M6C are present in the microstructure. M6C and G-phase are located in contact with M23C6 carbides or are found in areas of second-ary M23C6 carbides. Possibly a small fraction of the G-phase particles are rather M(C,N) particles which can, however, not be distinguished due to their small dimensions. In the outer sub-surface zone some M2(C,N) is present.

��30°F/�0�,30� h/�.�� ksi (��°C/�0�30� h/�4.� MPa)The microstructure of this part of the tube which had been exposed

to a slightly higher temperature is significantly different from that of the section exposed to 1607°F. (875°C) The primary and secondary carbides have changed their shape and have coarsened and the number of secondary carbides has decreased due to agglomeration, Figures 3-SE-Aa to 3-SE-Ar. The section contains besides 9.2% M23C6 also about 0.4% G-phase, 0.8% M6C and 0.6% M(C,N) in the core. In the metal sub-surface zone some M2(C,N) is observed, however no G-phase could be identified, Figures 3-SE-Ak to 3-SE-Ar. The different microstructures of both samples, A and B, cannot be explained by differences in compo-sition as their compositions are almost identical, Table 8.

��50°F/�4,000 h/0.�� ksi (��°C/�4000 h/4.�� MPa)The microstructure of this specimen which was exposed for 8.5

years reveals a similar appearance to that of the specimen exposed for 108,306 h at 1630°F. (888°C) Both show except of a higher Si content nearly identical compositions, Table 8.The structure is characterized by a eutectic-type of precipitate arrangement in areas of residual melt and along grain boundaries, Figures 3-SE-Ca to 3-SE-Cr. Besides about 8% M23C6 primary and secondary carbides, 1% G-phase and 2.2% M6C were measured. The M6C particles are sometimes found in combination with M23C6carbides in the dendrites of the matrix.

��50°F/�4��0 h/�.�� ksi (��°C/�4��0 h/�.0 MPa)After seventeen years of exposure the microstructure contains a large

number of creep cracks resulting from creep pores which had grown

together, Figures 3-SE-Da and 3-SE-Db. The M23C6 primary carbides have coarsened and secondary carbides have not been found. In the core area about 5.5% M23C6 carbides are present besides 1.6% M(C,N), Figures 3-SE-Da to 3-SE-Dj. The composition of this sample shows higher Si and Nb contents. In the sub-surface zone 6.2% M6C are found together with 4.6% M23C6 and the M(C,N) particles had been almost completely dissolved in M6C (0.2%), Figures 3-SE-Dk to 3-SE-Dr.

��°F/�3�,000 h/�.�� ksi (�05°C/�3�000 h/�3.� MPa)After about fifteen years of exposure at 1661°F (905°C) this speci-

men contains about 7.4% M23C6 primary and secondary carbides of a coarse structure. Furthermore in the core area as well as in the metal sub-surface zone about 2.0% M6C, 0.5% M2(C,N) and 0.1% M(C,N) are present, Figures 3-SE-Ha to 3-SE-Hr. The G-phase was not observed although the composition of the sample shows high Si and Nb contents. Possibly the high N content leads to bonding of Nb by the formation of M2N. In the metal sub-surface zone increasingly Nb and Cr-rich M2(C,N) are present.

��°F/�43,000 h/�.�� ksi (�05°C/�43000 h/�3.� MPa)After slightly longer exposure of sixteen years under comparable

conditions this specimen contains 6.5% M23C6 primary and secondary carbides, 2.2% M6C and 1.5% G-phase in the core, Figures 3- SE-Ia- to 3-SE-Ij. M(C,N) could not be detected. The composition of specimen I shows high Si and Nb contents. Although the N content is high, no M2N was formed. In the sub-surface zone no G-phase was present. The fraction of M23C6 primary and secondary carbides amounts to 7.9% and that of M6C to 6.2%, Figures 3-SE-Ik to 3-SE-Ir.

��°F/�43,000 h/�.�� ksi (�05°C/�43000 h/�3.� MPa)This specimen originates from the same tube as the two foregoing

specimens (H and I), however from a higher position. No differences in the microstructure were found compared to the specimen exposed for the same duration and same temperature shown before, Figures 3-

3�Atlas of Microstructures

SE-Ja to 3-SE-Jr. Moreover the volume fractions of the precipitates, as determined by image analysis, are almost the same as for the foregoing specimen. The chemical composition shows nearly identical contents, as well, Table 9.

�350°F/35,000 h/�.3 ksi (��°C/35000 h/�.�� MPa)The microstructure of this specimen which had been exposed to

the highest temperature over a duration of four years contains a high number of interdendritically oriented cracks which reach almost through the entire wall thickness. The precipitates are about 4% M23C6, mostly of the primary type, and 1.3% M(C,N) which was rich in Nb and also contained some Ti, Figures 3-SE-Ea to 3-SE-Ej. Locally clear indica-tions of beginning melting was observed. In these spots a fine lamellar eutectic M23C6 structure was present which partially contained some M6C, Figures 3-SE-Ek to 3-SE-Er. In the outer sub-surface zone around the M23C6 particles and within the dendritic matrix some local deple-tion features were observed which resulted in local over-etching of the microstructure, Figures 3-SE-Es to 3-SE-Ez.

ConclusionsAmong the two service-exposed tubes of the alloy HPMA only the

one exposed at 1800°F (982°C) for 7.5 years contained, besides M23C6 and residues of undissolved M(C,N), some G-phase in an amount less than 2%. The tube which had been exposed for eleven years at the lower temperature of 1673°F (912°C) showed M6C and M2(C,N) instead of G-phase. The chemistry of this specimen shows a very high N content which obviously leads to the formation of Nb-rich M2N. G-Phase cannot be formed because not enough free Nb is present. Possibly the highly Nb-containing M2(C,N) represents a precursor stage of the formation of the G-phase. The HP Modified service-exposed tubes with the exception of the two tubes with exposure time higher than 75,000 h contain between 0.4 and 1.6% G-phase besides 7-9% M23C6 and 0.8-2.5% M6C. In microstructures containing more than about 1% G-phase

no M(C,N) can be observed. One tube peaks out due to its high amount of M(C,N), while the amount of M23C6 is low. In this tube no G-phase had been formed which may be due to its chemical composition. Indeed a high content of Nb besides a high Si content was found which however does not explain why the G-phase had not formed here. Another tube contains Nb-rich M2(C,N) instead of G-phase. Again a high N content was found by chemical analysis. It can be suspected that the forma-tion of Nb-rich M2N is enhanced by high N contents. Again it can be suspected that M2(C,N) is a precursor stage of the formation of G-phase. In other words M2(C,N) is possibly dissolved after longer exposure times and G-phase is formed. Evidently at 2350°F (1288°C) a crit- ical temperature has been reached at which the material HP Modified shows molten areas in the microstructure due to the formation of low-melting phases which can reduce the creep strengths locally to a dramatic extent.

SummaryComposition of the Precipitates

In Table 13 the average compositions in atomic-% of the differ-ent types of carbides and phases are listed for the different alloys. The values of all elements except for C, (N), which were taken from the difference of the sum of the determined metal fractions and 100%, rest upon several averaged spot measurements for all samples of each alloy by EPMA. Comparison of the results shows a good agreement within the accuracy of the measurements for each type of phase. The existence of M2(C,N) of different compositions in alloy HPMA und HPNb is remarkable. Both alloys show probably Cr-Nitrides and Cr- and Nb-rich M2(C,N).

Precipitation Kinetics DiagramsThe development of the contents (area-%) of the different types of

carbides and phases as a function of exposure time on a logarithmic

3� Atlas of Microstructures

Table 13: Composition of the precipitates in atomic-% (Mean values of several measurements)Type Alloy Cr Ni Fe Si Nb Ti Zr C, (N)

35/45 63.7 1.1 2.4 0.0 0.0 0.2 0.0 32.6M7C3 HPMA 58.9 1.1 7.0 0.0 0.2 0.0 0.0 32.8

HPNb 61.9 1.0 5.9 0.0 0.2 n.a. 0.0 30.635/45 4.0 2.1 1.0 0.0 49.0 3.5 0.0 40.1

M(C,N) 35/45 3.1 1.3 0.5 0.1 8.6 30.5 9.4 46.8HPMA 3.6 2.1 1.6 0.6 48.4 2.1 2.1 39.5HPNb 2.8 1.2 1.4 0.1 49.9 0.1-1.1 1.6 42.935/45 68.9 3.4 3.3 0.0 0.1 0.0 0.0 23.9

M23C6 HPMA 64.0 3.4 8.9 0.0 0.1 0.0 0.0 23.6HPNb 66.5 3.4 8.7 0.0 0.1 0.0 0.0 21.135/45 34.0 27.1 1.9 13.0 8.3 0.2 0.0 15.5

M6C HPMA 38.0 23.9 3,4 12.0 4.6 0.0 0.0 18.0HPNb 31.3 25.8 3.6 13.1 10.2 0.0 0.0 15.935/45 5.1 49.4 2.9 22.0 18.9 0.5 0.6 0.2

G-Phase HPMA 2.7 47.3 5.3 22.9 20.0 0.4 0.0 0.9HPNb 4.0 47.7 5.8 21.8 19.6 0.0 0.1 1.235/45 64.8 0.7 0.8 0.0 0.3 n.a. n.a. 33.4

HPMA 62.6 0.3 1.5 0.0 0.3 n.a. n.a. 35.0M2(C,N) HPMA 32.0 1.8 3.0 0.0 29.8 1.4 0.0 32.0

HPNb 63.0 0.5 1.9 0.0 0.3 n.a. 0.0 34.4HPNb 32.2 1.3 2.5 0.0 31.5 0.1 0.0 32.4

Annotations: n. a. - not analysed


scale for different temperature ranges are represented in the diagrams in Figures 1-PKD-a to 1-PKD-d, 2-PKD-a to 2-PKD-c and 3-PKD-a to 3-PKD-d. In the plots, which reveal trends for the increase or decrease of the contents of the different precipitates, error bars for the standard deviation are drawn. As these trends are rather complex and the diagrams are not easy to understand in a further step for each temperature range the upper and lower limits for each type of precipitate were plotted in the form of scatter bands. The results are given in Figures 1-PKD-e, 2-PKD-d and 3-PKD-e, where the scatter bands are hatched in different colors for different temperature ranges. These plots show the range of fractions of MC, M23C6, M6C and Gphase expected for aged alloys of the types 35/45, HPMA and HP Modified for different temperature ranges based on the results of the image analysis measurements.

Alloy 35Cr45NiThe development of the M23C6-Carbides is shown in Fig. 1-PKD-a.

The content of M23C6 increases the faster the higher the temperature. The content of M(C,N) strongly depends on the composition of the alloy, especially on the alloy content of the M(C,N) forming elements Nb and Ti. The plot of a tendency does not make much sense here, Figure 1-PKD-b. For the formation of carbides of the M6C-type a similar trend is observed, Figure 1-PKD-c. The area fraction is up to 10.7% after expo-sure times of 5,000-10,000 hours at higher temperatures (1900-2050°F, 1038-1121°C). Higher temperatures and exposure times >1,000 hours seem to promote the formation of M6C, which is rich in Cr, Si and Nb. The G-phase is observed in the temperature range of about 1800-2000°F (982°C-1093°C) after exposure times of 500 <T<10,000 hours, Figure 1-PKD-d. For the formation of the G-phase high Nb fractions are neces-sary. Due to diffusion processes the NbC dissolves, which is shown by a decreasing content of the M(C,N). In addition a large fraction of Si and Ni is needed for the formation of the G-Phase. The supply of Ni and Si results from the matrix. It is remarkable that the M6C-content is lower or M6C does not exist at all, in samples, in which the G-Phase is present

and vice versa, cf. Figures 1-PKD-c and 1-PKD-d. The scatter bands of the area fractions of the precipitates for the different temperature ranges are shown in Fig. 1-PKD-e. From these scatter bands the user can easily extract the expected area fraction of a certain precipitate (MC, M23C6, M6C and G-phase) for the specified temperature ranges. It must be pointed out, that the fraction of MC strongly depends on the alloy composition and on the formation of M6C or G-phase, respectively. It is noticeable that the MC-content is lower or MC does not exist at all, in samples in which M6C or the G-phase are present and the that M6C-content is lower or M6C does not exist at all, in samples, in which the G-Phase is present.

Alloy HPMAThe growth of the M23C6-fraction with increasing exposure time for

three temperature ranges is similar to the development of this type of carbides in alloy 35Cr45Ni, Figure 2-PKD-a. The scatter of the amount of M(C,N) is shown in Figure 2-PKD-b. In the actual case its fraction predominantly depends on the composition of the sample, because here only in one case G-Phase is present. Figure 2-PKD-c shows the change in the fraction of M6C for the different temperature ranges. No clear tendency can be derived. In comparison to alloy 35Cr45Ni less M6C (max. 1.7%) is present. Only one of the samples was containing the G-phase. The scatter bands of the area fractions of the precipitates for the different temperature ranges are shown in Fig. 2-PKD-d. As only one of the SE-samples was containing the G-phase, no scatter bands for the G-phase were drawn. Attention should be given to the different scale of the y-axis for the M6C-area fractions in comparison to alloy 35Cr45Ni (Fig.1-PKD-e).


Alloy HP ModifiedFor this alloy the content of M23C6-Carbides is more stable during

exposure time. No clear dependence on the temperature range is shown in Fig. 3-PKD-a and 3-PKD-e compared to 35Cr45Ni (Fig. 1-PKD-a and 1-PKD-e) and HPMA (2-PKD-a and 2-PKD-d). The maximum fraction of M23C6 is 11%. One sample shows a very low fraction because of dissolution processes due to a temperature near the melting point (red value (SE-E) in Fig. 3-PKD-a and the beginning of melting, marked with an arrow in Fig. 3-PKD-e). The distribution of the M(C,N) frac-tions shows two levels depending on the respective content of Nb. High Nb-contents ≥ 1.18 mass-% seem to result in high M(C,N) fractions (1.0-3.0 area-%). Low Nb contents (0.74-0.83 mass-%) seem to result in lower M(C,N) fractions (≤1.0 area -%). In Fig. 3-PKD-e two green hatched scatter bands representing M(C,N) are shown for the tempera-ture range of 1787-1800°F (970-982°C). The formation of M6C initially shows an increase in its fraction (max. 4%) up to an exposure time of about 10,000 hours. After long exposure times >50,000 hours this type of carbides partly seems to dissolve, while simultaneously in certain cases G-Phase (max. 2%) is formed (Fig. 3-PKD-c and 3-PKD-e). G-phase was only observed in the temperature range of 875-955°C for exposure times between 750- 150,000 hours (Fig. 3-PKD-d and blue hatched scatter band in Fig. 3-PKD-e). In this temperature range in addition to M6C the G-phase exists. For higher temperature ranges G-Phase does not exist at all whereas M6C is sometimes present (green and red scatter bands in Fig. 3-PKD-e).

The Project Team that reviewed this document was surprised at the rapid conversion of MC to M23C6 carbides and at the absence of G phase in some samples. Following receipt of the manuscript Dr. Schütze offered some additional insights. This discussion is in the Addendum.

Acknowledgement

The financial support by MTI and the very pleasant collaboration with Galen Hodge, MTI are gratefully acknowledged.

Literature1L. Dillinger and D. Roach, Metallurgy and Metallography of Reformer

Alloys, Battelle Columbus Laboratories, May 1, 19752Metallografie, 14. Edition, Editors H. Schumann and H. Oettel,

2005 Wiley-VCH Verlag GmbH Co. KGaA, Weinheim, ISBN 3-527-30679-X

3Atlas der Interferenzschichten-Metallographie, Editors H.- E. Buehler and H. P. Hougardy, 1979 Deutsche Gesellschaft für Materialkunde e.V, Frankfurt am Main, ISBN 3-88355-015-9

4I. Aydin, H. - E. Buehler, Contrasting the Carbides M23C6, M7C3 and M6C in High Temperature Resistant Cast and Wrought Alloys, Praktische Metallografie 17, 1980, no. 9, 441 – 446

5I. Aydin, H. - E. Buehler and A. Rahmel, Precipitation in the heat resistant nickel-base cast alloys G-NiCr 28 W and G-NiCr 50 Nb, Arch. Eisenhuettenwesen 54, 1983, no. 11, 461 – 466

6I. Aydin, H. - E. Buehler and A. Rahmel, Determination of opti-cal constants of carbides and carbonitrides in refractory steels, Microstructural Science 9, 1981, 55 – 63

7I. Aydin, H-E. Buehler, A. Rahmel, Observations concerning nitrida-tion of refractory austenitic materials in air and combustion gases, Werkstoffe und Korrosion 31, 1980, 675 – 682

8ASM Handbook, Volume 9, Metallography and Microstructures, 1998, ASM International, Materials Parc; OH 44073

9Dr. D. Flahaut, D. T. Pool, 35Cr45Ni Micro-alloy – Friend or Foe, presented at: ERTC Petrochemical Conference, Vienna, 2004

10John J. Hoffman, G.E. Gapinski, Properties and Microstructures of Outlet Manifold Components, Ammonia Technical Manual, 2002

11R. C. Ecob, R. C. Lobb and V. L. Kohler, The formation of G-phase


Addendum1. Etchant CookbookDr. Vogel’s pickle:Vogel’sche Sparbeize (German name) is a commercial pickle, named

after its inventor Dr. Otto Vogel (Patents: DRP. 309264, 346530 und 350771; Brit. P. 158768,

188713;….)Dr. Vogel’s Sparbeize is a chemical solution whose exact recipe is not

known for the public. It partially acts as an inhibitor.Known contents are: ≤ 20% H2SO4, ≤ 1% H3PO4, 40-50% 1-

methoxy-2-propanol, 5-7% nonylphenol-ethoxylate, 3-5% thio-uric acid. It can be purchased e.g. from the following supplier:

Buehler GmbHIn der Steele 2D-40599 Düsseldorfhttp://www.buehler-met.dehttp://www.buehler.comPurchase order number : 170014 quantity: 0,25l price: 26,00€Safety data sheet: www.buehler-met.de/pdf/msds/170014.pdfMost probably there will also be an American Buehler subsidiary

with the address given on the above “.com” web page.

V2A etchantThis is a traditional etchant dating back to the middle of the last

century which had been developed for the etching of austenitic stainless steels. We took the composition from the chapter “etchants of Ni and Ni-alloys: recipe m10” of the following German book.

Metallographisches, keramographisches und plastographisches Ätzen

6. überarbeitete Auflage[Metallographic, keramographic and plastographic etching. A

Cookbook.]Unter Mitarbeit von: Carle, Veronika.1994. 6. Auflage. XII , 320 Seiten, 23 Abbildungen, 21x15cm(Materialkundlich-Technische Reihe, Band 1)Publisher: Gebr. Borntraeger, BerlinISBN 3-443-23014-8 brosch., EUR 26.00Similar etchants may be described also in American textbooks for

metallography.

As a cooking guide you can use:Composition: take 100 ml distilled H2Oadd successively100 ml HCl 1.19 M,10 ml HNO3 1.40 M,0.3 ml of an inhibitor e.g. Dr. Vogel’s Sparbeizestir up the mixture and heat up to 50°CEtching: move the sample in the hot etchant for 25-30 sec by usinga pair of tweezers

2. Rapid conversion of M7C3 to M23C6We know from experience, that the conversion of M7C3 to M23C6

is rapid in other comparable alloys, too (see Table 2 and Fig. 12, Fig. 38 and Fig. 66). If you consider the Time-Temperature-Preciptation-Diagramm of the cast alloy 1.4857 whose composition is similar to HP-MA and HPNb, you see that for a temperature of 900°C the trans-formation of M7C3 to M23C6 is already completed after 20 hours of exposure, Fig. 38. For higher exposure temperatures the conversion is even faster.

So it is not surprising that we didn’t find M7C3 and M23C6 in coexistence, because the exposure times and temperatures are outside the the band of conversion. The reaction of the conversion is 2M7C3 + 9M→ M23C6.


Source: I. Aydin, Ausscheidungsverhalten einiger hitzebeständiger, chromhaltiger Gußlegierungen auf Eisen– und Nickelbasis, Dissertation, RWTH Aachen, 1982

3�Atlas of Microstructures

3. Precipitation of the G-PhaseThe G-phase is an intermetallic silicide of the type Ni16Nb6Si7.

Many of the intermetallic phases do not show a broad range of compo-sitions, some are even socalled “line”-phases. Presumably stoichiometry does not allow for major variations concerning the composition also in the case of the G-phase. Intermetallics precipitate if specific equilibrium conditions are fulfilled.

Indeed the G- phase was found most frequently in the 35/45 alloy in a temperature range of 982°C-1150°C. Possibly the higher Ni-content in comparison to the HP-MA and HP-Nb alloys beside Nb and Si favours the G-phase formation. In the HP-MA alloy it was only found after exposure at 982°C / 65664hours. This service exposed sample contains 1.3 wt-% Si and 0.98 wt-% Nb measured after exposure. For the HP-Nb alloy the G-phase is only present in the temperature range of 875-955°C. The Si content of these samples varies between 0.9-1.7 wt-% and the Nb content between 0.74-1.2 wt-%. The service exposed HP-Nb samples show G-phase in a temperature range of 875-905°C for long time exposure of 75000-143000 hours. The wall thickness of these samples is higher in comparison to the creep rupture

samples. Probably the diffusion of carbide and G-phase forming elements is much slower at this lower temperature and the diffusion path of Si to the surface of the sample is longer in samples of larger dimensions. Si may be present for a longer time for the formation of G-phase than in thinner samples where it may be consumed by oxidation at the surface.

For the G-phase formation the following informations are known from the literature:

• Ageing of Nb-containing castings results in the replacement of Cr7C3 carbides by M23C6, secondary precipitates of M23C6 and NbC, with the replacement of some of the pre-existing NbC by the G-Phase

• The initial volume fraction of NbC and its rate of transformation to G-Phase were found to increase with the Si content

• Silicon The higher the Si content, the higher is the upper limit tempera-ture of formation for G-phase, and the faster is the kinetics of formation

• Niobium The higher the content of Nb the higher is the content of G-phase. Only at very high contents of Nb (more than 1.9 wt-%) the G-phase may partially be formed directly from NbC

• Titanium Ti addition transforms NbC into NbTiC which is more stable and less easy to destabilise to form G-phase. Ti therefore inhibits G-phase formation. Maximum efficiency is 0.1% wt-% Ti

• The following chemical composition (wt-%) is preferred in order to limit Gphase formation:

C Si Ni NbMin 0.40 1.3 42 0.7Max 0.45 1.6 44 0.9

Source: Dr. D. Flahaut, D.T. Pool, 35Cr45Ni Micro-alloy – Friend or Foe, presented at: ERTC Petrochemical Conference, Vienna, 2004


Atlas of Microstructures

Documents

alloy hpma exposure

system conditions

serviceexposed specimens

dechema e

xxcreep rupture test

limitation liability

usath e results

description of test