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1 Advanced Materials Research Online: ISSN: , Vols , pp doi: / Trans Tech Publications, Switzerland Phase Transformation Of P91 Steels Upon Cooling After Short Term Overheating Above Ac 1 & Ac 3 Temperature Ng Guat Peng 1, a, Badrol Ahmad 2,b, Mohd Razali Muhamad 3,c Mohd Ahadlin 4,d 1,2 TNB Research Sdn. Bhd. No 1, Jalan Air Hitam, Kawasan Institusi Bangi, 43000, Kajang, Selangor, Malaysia. 3,4 University Technical Malaysia/Post Graduate Center. Hang Tuah Jaya, 76100, Durian Tunggal, Melaka, Malaysia. a guatpeng@tnbr.com.my, b badrol@tnbr.com.my, c mohdrazali@utem.edu.my, d ahadlin@utem.edu.my Keywords: Ferritic, martensitic, PWHT, Ac 1 ; Ac 3, phase transformation Abstract. Advanced ferritic steels containing 9 wt% Cr are widely used in the construction of supercritical and ultra supercritical boiler components. Grade 91 is one of the most common alloys used in this application. The microstructure of the as supplied 91 materials consists of a tempered martensite matrix, a fine dispersion of intergranular chromium rich M 23 C 6 precipitates and intragranular carbonitrides MX particles rich in V and Nb. This steel requires post weld heat treatment (PWHT) to produce a tempered microstructure after welding to develop excellent creep strength for high temperature service. Based on past experience, situations may arise whereby the components are subjected to an accidental overshoot in temperature during PWHT. The consequence is the formation of deleterious phases which will result in undesirable changes in material property. In this research, P91 base metal specimens were heated to various peak temperatures in a laboratory furnace. Heat treatment parameters, as practiced at site, were applied. Peak temperatures applied were below Ac 1, between Ac 1 and Ac 3, and above Ac 3. Hardness measurement demonstrated a significant reduction once the Ac 1 temperature was exceeded, due to the presence of soft α-ferrite matrix. As the temperature was increased towards Ac 3, newly transformed fresh martensite which is hard and brittle in nature would form the dominant matrix. The phase transformation and precipitate morphology changes were studied using optical microscopy and scanning electron microscopy techniques. Three factors were identified to determine the phase transformation: (1) the homogeneity level and amount of precipitates dissolved in austenitic matrix upon heating; (2) slow cooling rate that may shift the cooling curve to enter ferrite nose and (3) deviation in chemical composition. Introduction Modified 9Cr-1Mo ferritic steel is an important structural alloy for modern steam generator application in electric power industries. These advanced alloys have been widely used in the construction of supercritical and ultra supercritical boilers [1-2]. One of the most commonly used steels is classified as grade 91. During the last twenty years, grade 91 (9Cr-1Mo-V-Nb) has been developed and successfully used in high temperature service. This steel is designed by adding strong carbide/nitride forming elements such as Nb, V and N in plain 9Cr-1Mo steels for creep strength enhancement. P91 was originally developed by the Oak Ridge National Laboratory for the application in the fast breeder reactor [3]. The steels are governed by the American standards ASTM 213 and 335. These materials exhibit a fully tempered martensitic structure with finely dispersed carbides, nitrides and carbonitrides that provide optimal combination of high creep strength and toughness [4-5]. The greatest challenge of this steel, from the operation history in power plants, is the need to conduct a stringent temperature control during post weld heat treatment (PWHT). PWHT is recognized as tempering process to improve material toughness. The tempering temperature All rights reserved. No part of contents of this paper may be reproduced or transmitted in any form or by any means without the written permission of Trans Tech Publications, (ID: , Pennsylvania State University, University Park, USA-06/03/16,22:53:41)

2 Advanced Materials Research Vols established for 91 steels is in the range of C to C which is just below Ac 1 [3]. The past experience showed that incidents sometimes occurred where the metal temperature unintentionally overshoot beyond Ac 1 and even Ac 3 due to faulty thermocouples, power surges or negligence. Previous research has shown that exceeding Ac 1 temperature would cause micro phase transformation to soft α-ferrite and hard fresh martensite during cooling [6]. Base on codes of practice, tubes or pipes will be discarded or disposed when they are overheated during PWHT. Studies associated with re-use of overheated tubes / pipes by de-rating or limiting them to temporary service are still rare at present stage. During emergency or dilemma where repair could not be done immediately due to lack of spares or no allowance to outage extension, the re-use of overheated tubes/pipes on short term basis could be an alternative solution. Nevertheless, there are no sufficient research data to support the use or prove the reliability of degraded tubes/parts on short term service after temperature excursion above Ac 1 and Ac 3. The changes in high temperature creep behavior after the short term damage are unclear and uncertain. Recent work [7] has shown that two factors are required to significantly change the microstructure of welded P 91 sections: (i) The peak temperature exceeds Ac 1 during PWHT and (ii) Slow cooling rate during PWHT. The present work is aimed to investigate the microstructure evolution of 91 base metal as a result of short thermal exposure above Ac 1 and Ac 3. Laboratory aging method was used to produce microstructure transformation at different peak temperatures above Ac 1 and Ac 3. Metallurgical Properties of Modified 9Cr-1Mo Steels The materials used in this research were 9 wt % Cr steel alloys designated as grade 91 in the form of pipe sections, hence subsequently designated as P91. Chemical composition, heat treatment and creep rupture strength of P91 steels are given in Table 1 [8]. The microstructure of the steel is classified as tempered martensite; supplied in normalized & tempered condition. Tempering in the range of C C leads to recovery of ductility by annihilation of dislocations and formation of ferrite sub grains [8]. Another important objective of tempering is the intention to develop precipitation strengthening mechanism, mainly derives from the dispersion of M 23 C 6 where M is (Cr,Fe,Mo) and fine Nb & V carbonitrides (MX) along prior austenite grain boundaries, ferrite subgrain boundaries and dislocation network. Table 1 : Typical chemical composition, heat treatment and mean creep rupture strength of steam pipe P91 steels Mass % P 91 C 0.09 Si 0.29 Mn 0.35 Cr 8.70 Mo 0.90 W - Ni 0.28 V 0.22 Nb N B - Austenisation C Tempering C 10 5 h creep rupture strength at C 94 MPa

3 1758 Advances in Chemical, Material and Metallurgical Engineering Experimental Procedure Sectioning of specimens The P91 specimens were cut from a 254 mm pipe with a wall thickness of 25.4 mm. The specimen size was about 100 mm in width, 25 mm in thickness and 30 mm in length, as shown in Fig. 1. The transverse sections of the specimens were extracted for metallographic examination after the laboratory aging heat treatment. The remaining parts were machined into a standard, dumb bell cylindrical creep specimens with threaded ends for creep rupture test. Virgin specimens of P91 were also prepared for metallographic examination as a baseline reference. Laboratory heat treatment The P91 specimens were heat-treated using a Nabertherm laboratory furnace. Certified type N thermocouple wires were mounted onto the specimens through a drilled hole of 1 mm diameter. The metal temperatures were transmitted, recorded and displayed by a calibrated data logger. The main purpose of this investigation was to examine the microstructure transformation resulted from a possible overshoot above the normal tempering temperature for this steel. The specimens were heated up, soaked for two hours at various temperatures and allowed to cool down to room temperature. The peak temperatures selected were just below Ac 1, between Ac 1 and Ac 3 and above Ac 3, covering the range between C and C in 50 0 C steps. The heat treatment curves are presented in Fig. 2. The selected temperatures are shown in Table 2. The heating rates used for P91 was C / hour. The cooling rates were controlled manually to be as close as to the heating rates until the temperature dropped to below C. Below this temperature, the cooling became unrestricted and the specimens were cooled continuously in still air. Optical microscopy, scanning electron microscopy and hardness measurement The laboratory aged specimens were subjected to metallographic examination and creep rupture testing. The metallographic specimens were mounted in resin, ground, polished and etched by immersion in Villella s reagent (1g picric acid, 5ml hydrochloric acid and 100 ml ethanol) for approximately 30s. The microstructure was revealed and observed using optical microscope and FE-SEM. Micro indentations were made on the specimens to measure the changes in hardness property using a test load of 0.3 kg in Vicker scale. The creep testing is still in progress and the results will not be presented in this paper. Table 2 : Peak temperatures used for P91 specimens Materials P91 Peak temperature, 0 C Fig. 1 P 91 specimens subjected to laboratory aging. Fig. 2 Heat treatment curve of P 91 specimen.

4 Advanced Materials Research Vols Experimental Results & Discussions Optical microscopy & scanning electron microscopy examination The as-received microstructure for P91 steel was observed to compose primarily tempered martensite with a uniform dispersion of fine precipitates. Specimens heated to C and C followed by cooling transformation contained a similar phase of tempered lath martensite. The precipitate morphology remained unchanged. From the SEM imaging, the as received, C and C specimens appeared to consist of a uniformly dispersed fine precipitates with aligned orientation. The precipitates were segregated along grain boundaries and dislocation network. (Fig. 3, Fig. 4 and Fig. 5). Fig. 3 As received microstructure. (a) optical micrograph; (b) SEM micrograph. Fig. 4 Transformed microstructure at room temperature after soaking at C. (a) optical micrograph; (b) SEM micrograph. Fig. 5 Transformed microstructure at room temperature after soaking at C. (a) optical micrograph; (b) SEM micrograph.

5 1760 Advances in Chemical, Material and Metallurgical Engineering It is normal that no visible transformation was observed at room temperature after exposure up to C as it was marginal and below Ac 1. The values quoted for Ac 1 for both steels are approximately C C, whereas Ac 3 is defined to be around C C [3]. The EDX analysis showed that majority of the precipitates were found to be rich in Cr, Fe and C. The precipitate composition suggested that they were mainly M 23 C 6 type, segregated along prior austenite grain boundaries (PAGB) and martensitic laths, as shown in Fig. 6. The measurement on the interior of the grain showed matrix composition. Fig. 6 EDX results showing the composition of precipitates in P91 as received specimen. Cooling transformation after temperature exposure at C had produced a change in the microstructure and precipitate morphology. The matrix had transformed to ferrite grains. There were evidence of precipitate coarsening and grain growth (Fig. 7). The microstructure appeared to be dislocation free. Fig. 7 Transformed microstructure at room temperature after soaking at C. (a) optical micrograph; (b) SEM micrograph.

6 Advanced Materials Research Vols Heating at C, followed by controlled cooling, the transformed microstructure was found to also consist of ferrite phase with precipitates. The population of the precipitates had reduced significantly. The precipitates appeared coarser. The ferrite grains were large. The distribution of the precipitates was observed to be random in nature and there was no specific alignment as shown in the as-received specimen (Fig. 8). Fig. 8 Transformed microstructure at room temperature after soaking at C. (a) optical micrograph; (b) SEM micrograph. The final transformed phase at room temperature, cooling from C and C are expected to produce untempered fresh martensite, as both temperatures had exceeded Ac 1 but below Ac 3. Both specimens should have transformed partially to -austenite phases which would form hard martensite at room temperature. Such transformation was not observed. The experiments at both temperatures were repeated. The heating condition remained unchanged but the cooling rates were raised by water quenching. In both cases, the quenched microstructures produced untempered martensite with a hardness in excess of 400 HV. The amount of precipitates had reduced significantly, as some of them had dissolved in austenite matrix during heating. The extent of dissolution depended on the peak temperature. Quenching from C produced a mixed microstructure of untempered martensite and ferrite with precipitation, which characterizes inter critical transformation between Ac 1 and Ac 3 (Fig. 9). Quenching from C produced an almost full untempered martensite with some precipitation, which represents upper critical transformation above Ac 3 (Fig. 10). This suggests that the lack of martensitic transformation in both temperatures with controlled cooling rates may be caused by the passage of cooling curves into ferrite nose. -austenite would have decomposed to soft α-ferrite before M s, the start temperature of martensite transformation at about C [9].

7 1762 Advances in Chemical, Material and Metallurgical Engineering Fig. 9 Quenched microstructure after soaking 2 hours at C, showing the presence of untempered fresh martensite (dark constituents) and ferrite (light constituents) Fig. 10 Transformed microstructure at room temperature after soaking 2 hours at C, showing an almost full transformation of untempered martensite. Previous research work suggested that during cooling, the austenite containing higher amounts of undissolved precipitates (carbides and carbonitrides) is less stable and would decompose to ferrite, while higher alloyed homogeneous austenite would form fresh martensite [6]. At C, a large amount of precipitates might have still remained as undissolved and the tendency to ferrite formation is high. At C, more precipitates, especially M 23 C 6 would have dissolved as solid solution in austenite matrix and smaller amount re-precipitated out during slow cooling, causing a reduction in the quantity. The transformed microstructure at cooling from temperature C appeared to contain a mixed phases in which there was evidence of an untempered martensite structure, co-existing with isolated ferrite grains. SEM image showed that the untempered martensite was in the shape of needle plates. Large ferrite grains were observed. (Fig. 11). Fig. 11 Transformed microstructure at room temperature after soaking at C. (a) optical micrograph; (b) SEM micrograph. When the specimen heated up to C was subjected to water quenching, it was found that the microstructure was transformed fully to untempered martensite with little precipitation (Fig. 12). At C which had exceeded Ac 3, the lack of complete martensite transformation at controlled cooling rates might be again caused by passage of cooling curve into ferrite nose due to lower speed of cooling.

8 Advanced Materials Research Vols Fig. 12 Transformed microstructure at room temperature after soaking 2 hours at C, showing a full transformation to untempered martensite. Microstructure transformed as a result of cooling from C at controlled cooling state had exhibited an untempered martensitic structure with little evidence of precipitates (Fig. 13). Cooling from C with slow cooling rate did not appear to be affected by ferrite formation. At C which is very close to the fully austenitizing temperature, most of the precipitates could have dissolved in austenite matrix and hence the austenite phase was stable, martensite was formed at the onset of M s, even at slower cooling rate. Fig. 13 Transformed microstructure at room temperature after soaking at C. (a) optical micrograph; (b) SEM micrograph, the dark patches are artifacts. Micro hardness measurement The hardness data following the laboratory aging thermal cycles at different peak temperatures between C and C for P91 specimens are presented in Fig. 14. The virgin, as-received specimens are of martensitic structure in tempered condition with a hardness 210HV-227 HV. The hardness remained unchanged up to C heat treatment. The hardness properties began to drop significantly at C and C cycles. The hardness had reduced to 165HV-187HV. This dramatic reduction in hardness was consistent with the microstructure changes to ferritic matrix with coarsening precipitates. A mixture of ferrite and untempered martensite at C produced an increase in hardness to the range of 252HV-367 HV. Upon cooling from C, fresh and untempered martensite was transformed to almost full extent and this explains the observed increase in hardness values to the range in excess of 400HV.

9 1764 Advances in Chemical, Material and Metallurgical Engineering Conclusions The results of this study demonstrate that exceeding the Ac 1 temperature during PWHT of P91 steels would have resulted in the formation of deleterious phases, for example, soft α-ferrite which has a poor creep strength and hard martensite which has a low toughness. The short excursion to high temperature between Ac 1 and Ac 3 would have also altered the precipitate morphology and population by dissolution and coarsening mechanisms. This is accompanied by a significant drop in hardness property if α ferrite formed the dominant matrix for peak temperature of C and C. The degraded specimens with soft α ferrite are expected to have a limited creep rupture life in high temperature service. As the peak temperature increased to C and C, hard and brittle martensite was formed, the transformation was in agreement with the prediction of CCT at high temperature above Ac 3. The presence of hard martensite without much precipitation is expected to have a low creep life as well, in addition to the risk of brittle cracking via mechanisms such as hydrogen induced cracking and etc. For heating above Ac 1, for example between C and C, quenched microstructure and as cooled microstructure at controlled rate C/hour contained different phases. Three factors were identified to govern the phase transformation. The first factor is associated with the amount of precipitates dissolved in austenitic matrix. At lower temperature after passing Ac 1, the M 23 C 6 particles are unable to fully dissolve; austenite is less stable and more prone to form α-ferrite [6]. It may be associated with a shift in ferrite nose during the continuous cooling transformation. This factor was less pronounced at higher temperature above Ac 3, such as C and C. The second factor is corresponded to cooling rates. At slower cooling rate, the cooling curves could have passed through ferrite nose, the unstable austenite would have decomposed to α ferrite and no martensite will be formed even after the temperature dropped to M s ( C). Another contributory factor may be the deviation in the chemical composition range of the specimens with respect to the chemical analysis for CCT diagram. Acknowledgements Fig. 14 Hardness profile of P91 specimens subjected to different peak temperature in laboratory aging. This work was accomplished in collaboration with Korea Electric Power Research Institute (KEPRI). The fund sharing and testing support from Green Energy Laboratory of KEPRI is gratefully appreciated.

10 Advanced Materials Research Vols References [1] B.K. Choudhary, E. Issac Samuel. Creep Behavior of Modified 9Cr-1Mo Ferritic Steel. Journal of Nuclear Materials. 412 (2011). Pp [2] S. Spigarelli, E. Quadrini. Analysis of the Creep Bahaviour of Modified P91 (9Cr-1Mo-NbV) weld. Materials and Design 23 (2002). pp [3] Nippon Steel Corporation, Data Package for NF616 Ferritic Steel, January 1993, second edition, March (1994). [4] A. Czyrska-Filemonowicz, A. Zielinska_Lipiec, P.J. Ennis. Modified 9% Cr Steels for advanced power generation: microstructure and properties. Journal of achievements in materials and manufacturing engineering, Vol 19, Issue 2, Dec [5] R.P. Chen, H. Ghassemi Armaki, K. Maruyama, Y. Minami, M. Igarashi. Microstructure Degradation during High Temperature Exposure Up to 10 5 H and Its Effects on Creep of Gr. 91 Steel. Proceedings from the sixth international conference on Advances In Materials Technology for Fossil Power Plants. August 31-September 3, 2010, Sante Fe, New Mexico, USA. Electric Power Research Institute, distributed by ASM International. pp [6] Dr. Alexandrov B., Wang L., Siefert J., Tatman J., Dr. Lippold J. Phase Transformations In Creep Strength Enhanced Ferritic Steel Welds. [7] J. Siefert, J. Sanders, J. Tanzosh, B. Alexandrov, J. Lippold, An Update of Phase Transformation During PWHT of Grade 91. Materials Science and Technology (MS & T) 2009, October 25-29, 2009, Pittsburgh, Pennsylvania. Joining of Advanced and Specialty Materials 2009 ( JASM XI). [8] John HALD and Leona KORCAKOVA. Precipitate Stability in Creep Resistant Ferritic Steels Experimental Investigation and Modeling. ISIJ International, Vol 43 (2003), No 3, pp [9] R.C. MacLachlan, J.J. Sanchez-Hanton and R.C. Thomson. The effect of simulated post weld heat treatment temperature overshoot on micro structural evolution in P91 and P92 power plant steels. Proceedings from the sixth international conference on Advances In Materials Technology for Fossil Power Plants. August 31-September 3, 2010, Sante Fe, New Mexico, USA. Electric Power Research Institute, distributed by ASM International. pp

11 Advances in Chemical, Material and Metallurgical Engineering / Phase Transformation of P91 Steels upon Cooling after Short Term Overheating above Ac 1 and Ac 3 Temperature /