Scholarly article on topic 'Development of Improved Materials for Structural Components of Sodium-Cooled Fast Reactors'

Development of Improved Materials for Structural Components of Sodium-Cooled Fast Reactors Academic research paper on "Materials engineering"

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{"Sodium-cooled fast reactors" / "austenitic stainless steel" / Nitrogen-enhanced / "Reactor structural" / "Ferritic steel" / Boron-added / "Type-IV cracking" / "Steam generator"}

Abstract of research paper on Materials engineering, author of scientific article — A.K. Bhaduri, K. Laha

Abstract Austenitic stainless steels (SS) with appropriate mechanical and irradiation-resistant properties are crucial for realizing sodium-cooled fast reactors (SFRs). In PFBR, with 40-years design life, 316LN SS with 0.02–0.03 wt-% C and 0.06–0.08 wt-% N, is used for structural components operating above 700K. For further improving high-temperature mechanical properties of 316LN SS, various heats with different nitrogen contents of 0.07, 0.14 and 0.22 wt-% were investigated. Detailed evaluation of tensile, creep, low cycle fatigue, creep-fatigue interaction and weldability, led to development of nitrogen-enhanced 316LN SS with 0.12–0.14 wt-% N having better combination of mechanical properties. Matching composition electrodes with 0.14 wt-% N have also been developed for its welding. The creep strength of fusion welded joints of modified 9Cr-1Mo (P91) steel, used for constructing the steam generators of PFBR, is considered to be a life limiting factor as a high percentage of the failures have been reported to occur in the inter-critical heat-affected zone, which is referred as Type-IV cracking. Micro-alloying the P91 steel with boron coupled with control of nitrogen content is considered as a possible way of reducing the strength disparities across the weld joint that enables minimising, if not eliminating, the Type-IV cracking problem. To optimize the boron and nitrogen contents, three heats of boron-added P91 steel, with 60-100ppm boron and 47-110ppm nitrogen, have been investigated. The weld joints of the boron-added P91 steel, with 60ppm boron and 110ppm nitrogen, exhibit better resistance to Type-IV cracking, with the creep rupture life of these weld joints being twice that of the weld joints of boron-free P91 steel.

Academic research paper on topic "Development of Improved Materials for Structural Components of Sodium-Cooled Fast Reactors"

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Procedía Engineering 130 (2015) 598 - 608

Procedía Engineering

www.elsevier.com/locate/procedia

14th International Conference on Pressure Vessel Technology

Development of Improved Materials for Structural Components of

Sodium-Cooled Fast Reactors

A.K. Bhaduria *, K. Lahaa

aIndira Gandhi Centre for Atomic Research, Kalpakkam 603102, India

Abstract

Austenitic stainless steels (SS) with appropriate mechanical and irradiation-resistant properties are crucial for realizing sodium-cooled fast reactors (SFRs). In PFBR, with 40-years design life, 316LN SS with 0.02-0.03 wt-% C and 0.06-0.08 wt-% N, is used for structural components operating above 700K. For further improving high-temperature mechanical properties of 316LN SS, various heats with different nitrogen contents of 0.07, 0.14 and 0.22 wt-% were investigated. Detailed evaluation of tensile, creep, low cycle fatigue, creep-fatigue interaction and weldability, led to development of nitrogen-enhanced 316LN SS with 0.12-0.14 wt-% N having better combination of mechanical properties. Matching composition electrodes with 0.14 wt-% N have also been developed for its welding. The creep strength of fusion welded joints of modified 9Cr-lMo (P91) steel, used for constructing the steam generators of PFBR, is considered to be a life limiting factor as a high percentage of the failures have been reported to occur in the inter-critical heat-affected zone, which is referred as Type-IV cracking. Micro-alloying the P91 steel with boron coupled with control of nitrogen content is considered as a possible way of reducing the strength disparities across the weld joint that enables minimising, if not eliminating, the Type-IV cracking problem. To optimize the boron and nitrogen contents, three heats of boron-added P91 steel, with 60-100 ppm boron and 47-110 ppm nitrogen, have been investigated. The weld joints of the boron-added P91 steel, with 60 ppm boron and 110 ppm nitrogen, exhibit better resistance to Type-IV cracking, with the creep rupture life of these weld joints being twice that of the weld joints of boron-free P91 steel. ©2015PublishedbyElsevierLtd. Thisisanopenaccess article under the CC BY-NC-ND license (http://creativecommons.Org/licenses/by-nc-nd/4.0/). Peer-review under responsibility of the organizing committee of ICPVT-14

Keywords: Sodium-cooled fast reactors; austenitic stainless steel; Nitrogen-enhanced; Reactor structural; Ferritic steel; Boron-added; Type-IV cracking; Steam generator

* Corresponding author. Tel.: +91-44-27480118; fax: +91-44-27480075. E-mail address: bhaduri@igcar.gov.in

1877-7058 © 2015 Published by Elsevier Ltd. This is an open access article under the CC BY-NC-ND license

(http://creativecommons.org/licenses/by-nc-nd/4.0/).

Peer-review under responsibility of the organizing committee of ICPVT-14

doi: 10.1016/j.proeng.2015.12.276

1. Introduction

Sodium-cooled fast reactor (SFR) technology represents the second of the three-stage nuclear programme envisioned for India, to make use of the large thorium reserves available [1]. In India, the Fast Breeder Test Reactor (FBTR) of 40 MW (thermal) has been operating successfully for over 25 years at the Indira Gandhi Centre for Atomic Research, Kalpakkam [1]. Based on this experience, a 500 MWe Prototype Fast Breeder Reactor (PFBR) has been designed indigenously and is at an advanced stage of construction, even as the design is being further optimised for enhanced economy with respect to cost of electricity production, for use in future reactors. All these components of SFRs are designed based on the ASME Boiler and Pressure Vessels Code and the French RCC-MR design code.

Austenitic stainless steels (SS) with appropriate mechanical properties are crucial for realization of SFRs. In SFRs, austenitic 316 SS and its variants are the preferred candidate materials for high-temperature reactor structural components operating above 700K, due to their adequate high-temperature tensile and creep strengths, compatibility with liquid sodium coolant, ease of fabrication, weldability and commercial availability. 316 SS used in SFRs differs from conventional grade of 316 SS w.r.t. close control on composition to avoid scatter in mechanical properties. In general, austenitic SS have relatively poor resistance to intergranular stress-corrosion cracking (IGSCC) in chloride and caustic environments. Weld joints of 316 SS exposed to chloride environments fail by IGSCC in the heat-affected zone (HAZ) due to combined influence of sensitization and presence of residual stresses introduced during welding. Hence, a nitrogen-alloyed low-carbon version (0.03 wt-% C max.) of this steel, 316LN SS, is chosen for high-temperature reactor structural components of the Indian Prototype Fast Breeder Reactor (PFBR). For the 316LN SS used in PFBR, 0.06-0.08 wt-% nitrogen is specified to compensate for loss in solid-solution strengthening due to reduced carbon content, with addition of nitrogen substantially increasing its creep rupture life. The beneficial effects of nitrogen arise due to higher solubility of nitrogen in the matrix than carbon, lower stacking fault energy of the matrix and introduction of strong elastic distortions in the crystal lattice, giving rise to strong solid solution hardening. Nitrogen also affects diffusivity of chromium in austenitic SS leading to retarding of coarsening of M23C6 thereby retaining the beneficial effects of carbide precipitation to longer times. For further improving high-temperature mechanical properties of 316LN SS, heats with different nitrogen contents were produced and evaluation of tensile, creep, low cycle fatigue, creep-fatigue interaction and weldability properties led to development of optimised nitrogen-enhanced 316LN SS with better combination of mechanical properties to enable achieving longer design life of 60 years or more. Matching composition welding consumables have also been developed for welding of the optimised composition nitrogen-enhanced 316LN SS.

Modified 9Cr-lMo (P91) steel is widely used in fossil power plants, petrochemical industries and many other heat transport systems, due to its good thermo-physical properties over austenitic stainless steels. P91 steel is being used for constructing the steam generators of PFBR. Compared to austenitic SS, the P91 steel has moderate creep strength coupled with high thermal conductivity, low thermal expansion coefficient and virtual immunity to stress corrosion cracking in chloride and aqueous media. However, with an increase in creep test temperature and decrease in applied stress, the weld joints of this steel, as well as those of other ferritic steels, fail in the fine-grain and/or inter-critical heat affected zone (FGHAZ / ICHAZ). This is due to lower creep rupture strength of the ICHAZ / FGHAZ than the other regions of the weld joint. Fracture of the weld joints in the FGHAZ / ICHAZ during creep test or in service is normally referred as Type-IV cracking [2-4]. To alleviate the problem of Type-IV cracking, the chemical composition of P91 steel can be altered by micro-alloying with boron coupled with control of the nitrogen content. Partial replacement of carbon in M23C6 type carbides with boron retards the coarsening kinetics of the carbides that contribute to reduction in the recovery of dislocations and consequent decrease in creep rate. Boron addition also influences the microstructural evolution in the HAZ during welding, and consequently the HAZ microstructure of boron-added ferritic steel weld joints is significantly different from that in conventional ferritic steels. Creep properties of boron-added ferritic steels indicate that the HAZ microstructure has significant beneficial effect on improving resistance of this class of steels to Type-IV cracking, as addition of boron to P91 steel produces a stable and relatively uniform microstructure across the HAZ. To accrue the beneficial effects of boron, nitrogen content has to be reduced substantially [5] to ensure that the brittle boronitride phase does not form. However, reduction in nitrogen content reduces the volume fraction of MX-type precipitates that are responsible for improving the high-temperature creep strength of P91 steel. Therefore, nitrogen content has to be optimised to ensure increased

volume fraction of MX-type precipitates without the formation of boronitrides. Hence, to optimize the boron and nitrogen contents three heats ofboron-added P91 steel, with 60-100 ppm boron and 47-110 ppm nitrogen, have been investigated.

This paper describes the R&D work carried out for developing these improved materials for application in SFRs.

2. Nitrogen-Enhanced 316LN Stainless Steel for Reactor Structural Components

To optimize the enhanced nitrogen content, the effect of nitrogen content on its tensile, creep, low cycle fatigue and weldability behavior has been investigated in four heats of 316LN SS with varying nitrogen content of 0.07, 0.11, 0.14 and 0.22 wt-% in which the rest of the composition was kept unaltered including the carbon content to 0.03 wt-% max.

2.1. Tensile Properties

In the temperature range 300-1123K, the yield stress and ultimate tensile strength increase linearly with increase in nitrogen content (Fig. 1) with the total elongation varying, as a function of nitrogen content, in the range of 4157% at 300K and 33-69% at 1123K [6].

Nitrogen, wt.% Nitrogen, wt.%

Fig. 1. Variation of (a) yield stress and (b) ultimate tensile strength with nitrogen content for nitrogen-enhanced 316LN SS.

2.2. Creep Properties

There is significant increase in creep rupture life with increase in nitrogen content at all stress levels in the range 140-225 MPa (Fig. 2), with the rupture life increasing by almost ten times with increasing nitrogen content from 0.07 to 0.22 wt-%. The beneficial effects of nitrogen arise due to higher solubility of nitrogen in the matrix than the carbon, reduction in stacking fault energy of the matrix and introduction of strong elastic distortions into the crystal lattice, giving rise to strong solid solution hardening [7]. Nitrogen also affects the diffusivity of chromium in austenitic stainless steels leading to retardation in coarsening of M23C6 thereby retaining the beneficial effects of carbide precipitation to longer times [8,9]. The increase in creep rupture strength with increasing nitrogen content could be correlated to decreasing tendency for sub-grain formation leading to uniform distribution of dislocations. However, creep rupture ductility decreases with increase in nitrogen content, with the variation of rupture ductility with rupture life showing a minimum in ductility at short durations and a continuous increase in ductility at longer rupture times, a trend characteristic of austenitic SS.

Nitrogen, wt.%

Fig. 2. Influence ofnitrogen on creep properties of 316LN SS at 923K.

2.3. Low Cycle Fatigue Properties

Studies on the influence of nitrogen content on low cycle fatigue (LCF) properties at various strain amplitudes between 300 and 873K [10] showed that the cyclic stress response curve is generally characterised by initial hardening, followed by saturation and then cyclic softening. The LCF life increases with increasing nitrogen content upto 0.14 wt-%, beyond which the fatigue life saturates/decreases (Fig. 3). The increase in fatigue life with increase in nitrogen content has been attributed to increasing planar glide of dislocations and better slip reversibility (i.e., less slip localisation), whereas reduction in fatigue life at nitrogen contents above 0.14 wt-% has been attributed to high matrix hardening and consequent decrease in residual ductility. Fatigue crack initiation and propagation is transgranular, with the crack initiation taking place along slip bands on the specimen surface. Initial crack propagation occurs over only one or two grain diameters along slip planes oriented at around 45° to the applied stress axis (stage I), followed by transition to stage II cracking that shows up as striations on the fracture surface. With respect to LCF properties, 0.14 wt-% nitrogen is the optimum content for the nitrogen-enhanced 316LN SS.

Nitrogen, wt. %

Fig. 3: Influence ofnitrogen content on LCF life of316LN SS atvarious temperatures.

2.4. Weldability and Welding Consumable

Evaluation of the hot cracking susceptibility of the different heats of 316LN SS with varying nitrogen content was carried out [11]. Gleeble-based hot ductility tests revealed the on-cooling hot ductility behaviour of nitrogen-enhanced 316LN SS. The effect of nitrogen content on nil strength, nil ductility and ductility recovery temperatures, NST, NDT and DRT, showed that, with increasing nitrogen content, NST decreases upto 0.14 wt-% N beyond which it increases again, while its effect on DRT is the reverse. However, NDT decreases linearly with increasing nitrogen content. In hot ductility tests, nil ductility range (NDR = NST - DRT), which measures how fast ductility recovers during on-cooling test, is 40K and 50K for 0.07 wt% N and 0.22 wt-% N steels, respectively, with the minimum being 30K for the 0.14 wt-% N steel. The Varestraint tests confirmed that the 0.22 wt-% N steel has the highest hot cracking susceptibility. Based on the weldability considerations, 0.14 wt-% N is found to be the optimum for the nitrogen-enhanced 316LN SS.

The 0.14 wt-% nitrogen in this optimised 316LN SS composition shifts the solidification mode of the weld metal to fully austenitic region, including that due to dilution of nitrogen from the base metal, thereby increasing its hot cracking susceptibility. This necessitated development and qualification of shielded metal arc welding electrodes, with chemical composition of weld metal specified to contain 0.14 wt-% N based on evaluation using the WRC-1992 diagram [12]. To ensure ferritic mode of solidification and sufficient delta-ferrite for ensuring adequate resistance to hot-cracking, the chemical composition of the weld metal for nitrogen-enhanced E316-15N electrodes was suitably tailored. Using the E316-15M welding electrodes developed for PFBR as the basis [13], for the E316-15N specification, the content of three elements C, N and Cr were adjusted within the ASME specification window (Table 1). Also, to ensure adequate resistance to both hot cracking and elevated-temperature formation of sigma-phase, delta-ferrite content of 3-7 FN is specified for the E316-15N weld metal.

Table 1: Chemical composition (in wt. %) and delta-ferrite content (in FN) ofE316-15weld metal for welding enhanced-nitrogen 316LN SS

Consumable C Cr Ni Mo N Mn Si P S Ti+Nb+Ta Cu Co B FN

E316-15 0.08 17- 11- 2.0- 0.5- 0.9 0.04 0.03 0.75

(ASME) max 20 14 3.0 2.5 max max max max

E316-15M 0.045- 18- 11- 1.9- 0.06- 1.2- 0.4- 0.025 0.020 0.1 0.5 0.2 0.002 3-

(PFBR) 0.055 19 12 2.2 0.10 1.8 0.7 max max max max max max 7

E316-15N 0.040- 19- 11- 1.9- 0.12- 1.2- 0.4- 0.025 0.020 0.1 0.5 0.2 0.002 3-

(New FBRs) 0.050 20 12 2.2 0.16 1.8 0.7 max max max max max max 7

Further, care has to be exercised during welding to avoid formation of nitrogen porosity. These modified E316-15N electrodes have been successfully developed, during which the most challenging was achieving the specified minimum toughness of 3.0 daJ.cm-2 after 1023K/100h ageing heat treatment. This requirement assesses the susceptibility of the weld metal to embrittlement by sigma-phase by transformation of delta-ferrite during high-temperature exposure, and for which the Mo content has to be carefully controlled to the lower limit. Another challenge has been improving the slag detachability of the deposited weld metal, as poor slag detachability often necessitates extensive grinding and rework. All these necessitated evaluation of a large number of trial batches for optimizing both the composition of the weld metal and the flux.

3. Boron-added Modified 9Cr-lMo Steel for Steam Generators

P91 steel, which is used in the normalized and tempered condition, has a tempered martensite structure, in which the additions of vanadium, niobium and nitrogen ensure intragranular precipitation of highly stable vanadium- and niobium-carbonitride (MX) particles on tempering and during creep exposure to confer relatively high creep strength. The detailed microstructure in the HAZ of ferritic steels is extremely complex and is controlled by the interaction of thermal fields, produced by the heat input from the welding process, and the phase transformation and grain growth characteristics of the materials being welded. Further modifications in microstructure can occur as a result of tempering either during the later stages of multipass welding and post-weld heat-treatment (PWHT) or

during service. These microstructures, which generally vary from wrought base material through transformed HAZ regions to cast weld metal, can have greatly different mechanical properties. Hence, fusion welded joints of P91 steel, as well as other ferritic steels, suffer from reduction in creep rupture strength compared to its base metal. The strength reduction is more at higher test temperatures and for longer creep exposures. The premature failure of the steel weld joint, termed as Type-IV failure, occurs in the softer FGHAZ / ICHAZ compared to the coarse grain HAZ (CGHAZ) and unaffected base metal. Microstructural modification of the base metal during weld thermal cycle and subsequent post-weld heat-treatment (PWHT) results in highly heterogeneous microstructure across the HAZ. The tri-axial state of stress due to creep strength inhomogeneity across the weld joint induces localized creep cavitation coupled with preferential creep deformation in the soft ICHAZ leading to Type-IV failure [2-5, 14-20]. The loss in mechanical strength of the ICHAZ of the weld joint is associated with coarsening of M23C6 type of carbides, recovery of martensitic-lath structure and formation of fine prior-austenite grain structure.

Therefore, fabricating weld joints with uniform microstructures having comparable mechanical properties across the joint would reduce triaxiality in the weld joint resulting in relatively uniform deformation and increased creep rupture life. In this direction, it has been already been shown that boron addition in P91 steel not only produces stable microstructures, it also produces relatively uniform microstructure across the HAZ of the weld joint [2-5]. The higher creep strength of boron-added steel is attributed to the stable microstructure due to the presence of fine M23C6 precipitates, which reduce recovery of dislocations and decreases the creep rate. However, there is a necessity to optimize the boron and nitrogen contents to use the beneficial effects of these elements by ensuring that the boronitride precipitates do not form due to marginal increase in their contents. Instead, the beneficial effect of nitrogen can be utilised through the formation of MX-type of precipitates, which further reduces the movement of free dislocation and further decrease in creep rate.

Boron-added P91 steel with varying boron and nitrogen is used for the present study, were received from an Indian manufacturer in the plate form. The chemical composition of the base metal used in this investigation is given in Table 2. From the table, it is evident that in addition to marginal differences chromium, nitrogen content is varied from 47 to 110 ppm while boron content is varied from 60 to 100 ppm.

Table 2: Chemical composition (wt. %) ofboron-added P91 steels; P91 specification ofPFBR and RCC-MR included for comparison

Alloy C Si Mn P S Cr Mo V Nb B N Ni Cu A1 Fe

P91BN1 0.09 0.35 0.40 0.004 0.005 9.0 1.00 0.20 0.07 0.010 0.0047 <0.02 <0.02 <0.005 Bal.

P91BN2 0.10 0.36 0.40 0.004 0.004 9.0 1.00 0.21 0.07 0.009 0.0100 <0.02 <0.02 <0.005 Bal.

P91BN2 0.10 0.48 0.50 0.004 0.004 9.2 1.00 0.21 0.07 0.006 0.0110 <0.02 <0.02 <0.005 Bal.

P91BN 0.08- 0.20- 0.30- 0.02 0.01 8.0- 0.85- 0.18- 0.06- 0.004- 0.004- 0.02 0.02 0.04 Bal.

(Specs.) 0.12 0.50 0.50 max max 9.5 1.05 0.25 0.10 0.011* 0.011 max max max

P91 0.08- 0.20- 0.30- 0.02 0.01 8.0- 0.85- 0.18- 0.06- NS 0.03- 0.40 0.04 Bal.

(PFBR) 0.12 0.50 0.60 max max 9.5 1.05 0.25 0.10 0.07 max max

P91 (RCC- MR) 0.080.12 0.200.50 0.300.50 0.02 max 0.01 max 8.09.0 0.851.05 0.180.25 0.060.10 NS 0.030.07 0.02 max 0.10 max 0.04 max Bal.

3.1. Microstructure of base metal

The microstructures of the normalized and tempered boron-free and boron-added base metals consist of tempered lath martensite having high dislocations density and precipitates decorating the boundaries. Comparison of the typical TEM bright field image of the boron-free and boron-added P91 steel base metals (Figs. 4(a) and 4(b) respectively) clearly demonstrate that the precipitates are finer in the boron-added base metal than in the boron-free base metal.

3.2. Microstructure of heat-affected zone

The microstructures of CGHAZ, FGHAZ and ICHAZ of boron-free and boron-added steel weldjoints are shown in Figs. 5(a) to 5(f). These photomicrographs were taken from the HAZ of weld joints at 0.5, 1.0 and 1.5 mm distance from the weld interface. The micrograph of the CGHAZ of boron-free steel shows clearly defined martensitic lath structure (Fig. 5a); but the lath structure is not clearly defined in the microstructure of both the FGHAZ (Fig. 5c) and ICHAZ (Fig. 5e). A large variation in prior-austenite grain size in different parts of the HAZ of the boron-free steel weld joint is evident. On the other hand, the lath structure is clear in the CGHAZ (Fig. 5b), FGHAZ (Fig. 5d) and ICHAZ (Fig. 5f) of the boron-added steel weld joint. Figure 5(f) shows very fine grains (marked by arrows) decorating the original prior austenite grain boundaries in the ICHAZ of the boron-added steel weld joint. Though, prior austenite grain size is similar in different regions of the HAZ in boron-added steel weld joint, boron-added, the different HAZ regions are referred to as CGHAZ, FGHAZ and ICHAZ based on peak temperature attained in these HAZ regions during the weld thermal cycle.

Fig.4. Typical bright field TEM images of (a) boron-free and (b) boron-added P91 steel base metals.

Fig. 5. Photomicrographs of different HAZ regions of (a, c & e) boron-free and (b, d& f) boron-added P91 steel weldjoints: (a & b) CGHAZ, (c & d) FGHAZ and (e &f) ICHAZ.

3.3. Hardness distribution across weld joint

The hardness distributions across the weld interface of the boron-free and boron-added P91 steel weld joints are shown in Fig. 6. The hardness of the weld metal is ~275 VHN in both the weld joints. The hardness of the CGHAZ at 0.5 mm from the weld interface of the boron-free and boron-added P91 steel weld joints is 267 and 235 VHN, respectively. The minimum hardness, at 2.0 mm from the weld interface in the ICHAZ, is 187 and 203 VHN, respectively for these two weld joints. This figure also shows that the hardness of the boron-added steel base metal is higher than that of the boron-free steel base metal. The hardness values are circled at two locations in the graph, one in the weld metal very close to the weld interface and the other in the ICHAZ. The former reveals decrease in hardness in the weld metal very close to the weld interface and the latter shows that the hardness of the ICHAZ is higher in the boron-added steel weld joint than in the boron-free steel weld joint. This drop in hardness in the weld metal of the boron-added steel weldjoint is due to presence of delta- ferrite.

u -|—i—i—i—i—i——i—i—i—i—i—i—i—i—i—i— 3 -2 -1 0 1 2 3 4 5

Distance across the weld interface, mm

Fig. 6. Hardness distribution across the weld interface ofboron-free (P91) and boron-added (P91B) steel weldjoints.

3.4. Tensile strength of weld joints

For boron-free and boron-added P91 steel weldjoints, the ultimate tensile strengths at room temperature are 659 and 634 MPa, respectively, while those at 823K are 433 and 474 MPa, respectively. Therefore, addition of boron does not lead to any significant difference in the tensile strengths of P91 steel weld joints at room and elevated temperatures. Also, for all the tensile tests, the location of failure is in the base metal close to the HAZ/base metal interface.

3.5. Creep properties

Studies have been carried out on 100-ppm boron-added standard (as per PFBR specification) P91 steel containing about 500-ppm nitrogen. Creep studies on this steel and its fusion welded joint have shown no beneficial effect of boron on creep rupture strength of the boron-added P91 steel containing 500-ppm of nitrogen and its weld joint (Fig. 7). This confirms that the contents of both boron and nitrogen in the P91 steel have to be judiciously controlled to accrue the benefits of boron addition by avoiding boronitride formation thereby ensuring that the free boron is adequately available.

Creep tests, carried out on three heats of P91 steels with controlled boron and nitrogen at 873K over different stress levels clearly show that the boron-added P91 steel, with 60 ppm boron and 110 ppm nitrogen, exhibits better creep rupture strength (Fig. 8).

175 150

tC 125 Q.

tn <n CD

Modified 9Cr-1Mo

0.01 wt. % Boron

0.05 wt. % Nitrogen

650 °C

■ P91-Base

• P91 -Boron-Base

A P91-Joint

Y P91-Boron-Joint Joint

10 100 1000 10000

Rupture life, hours

Fig. 7: Effect of 100-ppm boron on creep rupture strength ofP91 steel, containing 500-ppm nitrogen, and its weldjoint.

Rupture life, hour

Fig. 8. Variation ofcreep-rupture life at 873K and different applied stress forthe boron-free and the three boron-added P91 steel base metals.

■ P91 Base

• P91 Joint

y P91BJoint(1050°C/1h)

A P91BBase(1050°C/1h)

P91B joint(1150 °C/1h)

600 °c —» test in progress

10 100 1000 10000 Rupture Life, hours

Fig. 9. Variation ofcreep-rupture life at 873K and different applied stress for the boron-free and the three boron-added P91 steel weldjoints.

Fig. 10. Photographs oftransverse-weld specimens creep tested at 600 °C and 120 MPa for (a) boron-free and (b) boron-added P91 steel.

Creep tests, carried out as 873K at different stress levels clearly show that the weld joint of the boron-added P91 steel exhibits better resistance to Type-IV cracking, with the creep rupture life of these weld joint being twice that of the weld joints of boron-free P91 steel (Fig. 9).

Photographs of transverse-weld creep specimens of the boron-free and boron-added P91 steels tested at 873K and 120 MPa are shown in Figs. 10(a) and 10(b), respectively. The fracture location in the boron-free steel weld joint is close to the HAZ/base metal transition zone within the ICHAZ (Fig. 10a), which is the typical Type-IV cracking location. However, the fracture location in the boron-added steel weld joint is in the weld metal (Fig. 10b). These results show that there is not only an improvement in the rupture lives, but also there is a shift in the fracture location indicating that the boron-added P91 steel weld joint does not show typical Type-IV cracking behaviour at the temperature and stress level of testing.

4. Concluding Remarks

Detailed evaluation of tensile, creep, low-cycle fatigue, creep-fatigue interaction and weldability led to the development of an optimised nitrogen-enhanced 316LN SS with 0.12-0.14%N having better combination of mechanical properties.

Similarly, detailed studies on the effect of boron and nitrogen on the creep behaviour of modified 9Cr-lMo steel showed that P91 steel with 60 ppm boron and 110 ppm nitrogen exhibits remarkably improved creep strength with improved Type-IV cracking resistance compared to the conventional boron-free P91 steel.

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