characteristics and effects of austenite resulting from

7/31/2019 Characteristics and Effects of Austenite Resulting From

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Characteristics and effects of austenite resulting from

tempering of 13CrNiMo martensitic steel weld metals

P.D. Bilmesa,b,*, M. Solarib, C.L. Llorentea,c

aDepartamento de Mecanica, Facultad de Ingeniera, Universidad Nacional de La Plata, Calle 1 y 47, 1900 La Plata, ArgentinabConsejo Nacional de Investigaciones Cientficas y Tecnicas (CONICET), Av. Rivadavia 1917, 1033 Capital Federal, Argentina

cComision de Investigaciones Cientficas de la Provincia de Buenos Aires (CICPBA),

Calle 52, c/120 y 121, 1900 La Plata, Argentina

Received 31 July 2000; accepted 28 September 2000

Abstract

Low-carbon 13CrNiMo martensitic steels are remarkable for their high strength and high resistance to brittle

failure while retaining corrosion resistance together with weldability. These properties can be obtained when an

intercritical tempering is applied as heat treatment or postweld heat treatment (PWHT); promoting the

precipitation of finely distributed austenite that remains untransformed after cooling. The content and stability of

this austenite in the weld metal accounts for the high toughness even under subzero conditions. Transmission and

scanning electron microscopy (SEM), X-ray diffraction, and Mossbauer spectroscopy were used to study both the

austenite resulting from intercritical tempering of these soft martensitic stainless steel weld metals and the

austenitefracture interactions. To recognize the effect of the austenite content on impact toughness, single- and

two-stage tempering have been applied and evaluated through Charpy tests. The studies have shown the austenite

to be thermally stable, mainly due to its substructure, but not mechanically stable, indicating that the toughening

mechanism of the austenite particles is associated with transformation-induced plasticity (TRIP). D 2001 Elsevier

Science Inc. All rights reserved.

Keywords: Retained austenite; 13Cr-NiMo steels; Post weld heat treatments; Microstructure; Mechanical properties

1. Introduction

Soft martensitic stainless steels such as 13Cr

NiMo are widely used for hydraulic turbines, valve

bodies, pump bowls, compressor cones, impellers,

and high-pressure pipes in power generation, offshore

oil and gas, and petrochemical industries [1,2]. It is

known these steels perform well in applications

where corrosion and cavitation erosion resistance

are required. In addition, they have some resistance

to stress corrosion cracking (SCC) in CO2 and H2S

environments, high strength, and toughness even at

low temperatures or in thick cross-sections, and have

excellent weldability [36].

These alloys solidify to delta ferrite crystals [7].

The transformation of delta ferrite into austenite crys-

tals starts at around 1300C and ends, in the case ofequilibrium conditions, at around 1200C. With theactual cooling rates experienced during a welding

operation, small amounts of delta ferrite are super-

cooled during the delta ferrite) austenite transforma-

1044-5803/01/$ see front matterD 2001 Elsevier Science Inc. All rights reserved.

PII: S 1 0 4 4 - 5 8 0 3 ( 0 0 ) 0 0 0 9 9 - 1

* Corresponding author. Tel.: +54-221-423-6692; fax:

+54-221-425-9471.

E-mail address: [email protected] (P.D.

Bilmes).

Materials Characterization 46 (2001) 285 296


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tion. Similarly, the austenite)martensite transforma-tion leads to a martensitic microstructure containing

small quantities of supercooled retained austenite.

Thus, after cooling has been completed, the structure

consists of soft martensite with small amounts of

supercooled delta ferrite and austenite.

When a subsequent intercritical tempering at

600C (slightly above the Ac1 temperature) is applied

as postweld heat treatment (PWHT), excellent tough-

ness properties are usually obtained. This temperingpromotes the softening of the martensite and the

precipitation of finely distributed austenite along the

martensite interlath boundaries and prior austenite

grain boundaries. After the tempering, this austenite

remains untransformed and it is known that these

particles account for the high toughness of this alloy.

However, the thermal and mechanical stability of the

austenite particles and the mechanism by which the

austenite enhances the toughness properties greatly is

not yet well understood. Several theories [8,9] have

attempted to explain how this happens in a compositemicrostructure of tempered lath martensite with small

particles of austenite phase densely distributed along

the lath boundaries and the prior austenite grain

boundaries. Among these theories, the crack blunting

model suggests that a crack being propagated through

the steel would blunt in the ductile stable austenite

(fcc). Although some evidence of this mechanism

was presented in a previous work [10], recent results

show that the transformation of austenite (fcc) into

martensite (bcc) happens during the fracture process

[11]. These new results are in agreement with themodels that are based on the transformation of

austenite into martensite by a localized transforma-

tion-induced plasticity (TRIP).

The present work shows the chemical and sub-

structural characteristics of the austenite particles,

present in 13Cr NiMo weld metals after intercritical

tempering, and their thermal and mechanical stability.

The effect of both single- and two-stage tempering on

the austenite content of the weld metal, and thus on

the impact toughness, is studied. Additionally, the

mechanism by which austenite particles improve thetoughness of these materials is recognized.

2. Experiments

Automatic gas metal arc welding was used to

prepare the weld metals. Multiple-pass welds were

performed on AISI 410 plates using a 13Cr NiMo

welding wire. The composition of the welding wire

and the welding parameters are shown in Tables 1

and 2. The applied PWHT were a single-stage tem-

pering (intercritical tempering at 600C/2 h/air) and atwo-stage tempering (first at 670C/2 h/air with a

previous solution annealing at 950C/1 h/air, second

at 600C/2 h/air). Chemical composition, microstruc-

tures, tensile properties, and Charpy V-notch impact

energy were determined in the as-welded and tem-pered conditions. The chemical composition of the

weld metal was measured by an optical emission

technique (except for C, N, O, and S that were

measured by combustion analysis). A Philips 515

scanning electron microscope (SEM) operated at 20

kV was used to observe the microstructures of the

PWHT weld metals and the fracture surfaces of the

Charpy V-notch impact specimens. Samples were

ground and electropolished, the electrolyte composi-

tion for the latter phase being: 62 ml HClO4, 700 ml

ethanol, 100 ml butyl cellusolve, and 137 ml H2O.The specimens were then etched in Vilellas solution.

The Charpy V-notch impact tests were performed at

20C and 77C.The volume fraction of the austenite in the tem-

pered conditions were measured by X-ray diffraction

from a Rietveld analysis. The carbon concentration in

this phase was evaluated by X-ray diffraction through

the lattice parameter of the austenite. X-ray diffraction

patterns were obtained at room temperature with a

Philips PW1710 diffractometer, furnished with a

diffracted beam graphite monochromator. Data werecollected using CuKa radiation in the range

10 2q 120 at a step interval of 0.02. A Rietveldanalysis was performed using the program DBWS-

9411 [12]. The sample displacement, the background

(modeled with a fifth-degree polynomial), the unit cell,

the preferred orientation, the pseudo-Voigt profile

parameters, and the scale factor of the different phases

present in the sample were refined independently but

not simultaneously. From the Rietveld analysis, a

relative weight fraction was assigned to the refined

Table 1

Nominal composition of welding wire (wt.%)

Grade Wire C Si Mn Cr Ni Mo

ASME SFA-5.9 class ER410NiMo 1.6 mm 0.03 < 0.95 0.65/0.9 12/13.5 5.5 0.5

Table 2

Welding parameters

Gas

Gas flow

(l/min)

Voltage

(V)

Current

(A) DC +

Arc speed

(mm/min)

Interpass

temperature

(C)

Ar 20 25 255 150 < 120

P.D. Bilmes et al. / Materials Characterization 46 (2001) 285296286


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Fig. 1. Ductile fracture surface of Charpy-V notch specimen: (a) Low magnification 1000, magnification bar is 10 mm, (b) High

magnification 6000, magnification bar is 10 mm.

Fig. 2. X-ray diffraction patterns of austenite: (a) Single-stage tempering, (b) Two-stage tempering.



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ical stability of austenite particles was also evaluated in

regions close to the fracture surfaces of Charpy speci-

mens in the plastic zone preceding the crack front. This

was performed by means of X-ray diffraction through a

Rietveld analysis. Microhardness was measured in

these zones by using a Shimadzu Vickers meter with

a load of 20 g and a loading time of 15 s.

3. Results and discussion

The chemical composition of the weld metal in

each of the different conditions is given in Table 3.The compositions were similar in all conditions and

in line with the nominal composition of the consum-

able electrode. The nitrogen and oxygen contents are

similar for both PWHT conditions. Tensile properties,

Charpy V-notch impact energy, hardness, and auste-

nite contents of the weld metals before and after

PWHT are shown in Table 4.

The yield and tensile strength were higher for the

as-welded martensitic condition. Two-stage tempering

promoted a greater softening than single-stage tem-

pering. The highest values of impact strength, or

toughness, were observed following the two-stage

tempering. This is the condition that contains the

greatest amount of austenite. As regards the fracture

mode, it is worth mentioning that all PWHT speci-mens tested both at room and subzero temperatures

displayed 100% of fibrous fractures with typical

Fig. 3. Difraction pattern with (y iobs) and (y i

cal) for Rietveld analysis. Two-stage tempering condition.

Fig. 4. Difraction pattern of M2X for two-stage tempering condition.

P.D. Bilmes et al. / Materials Characterization 46 (2001) 285296 289


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dimples and without cleavage. Fig. 1a is representa-

tive of the fracture surfaces seen, this one from a

specimen tested at subzero temperature for single-

stage tempering. The typical dimple appearance asso-

ciated to ductile-dimple fracture with a high micro-

void density can be observed. This may be attributed

to the existence of the large number of internal

interfaces due to both nonmetallic inclusions and

austenite or transformed austenite particles, which

may act as void nucleation sites [14]. Fig. 1b shows

typical dimples with several small austenite or trans-formed austenite particles within them.

The X-ray diffraction patterns of the samples

corresponding to both single- and two-stage temper-

ing (Fig. 2a and b, respectively) were analyzed using

the DBWS-9411 program for Rietveld analysis. This

method was developed by Hugo Rietveld in 1969,

and today, it is one of the most powerful techniques

used for structural analysis, measuring reticular para-

meters, and other investigations about crystallogra-

phy. It consists of fitting step by step the experimental

intensities (yiobs) corresponding to all spectra, withthose (yi

cal) based on a specific crystalline structure

model, diffraction optical effects, instrument factors,

and other sample characteristics. The parameters

included in the model are refined to achieve the best

fitting of minimum squares of thousand of yi belong-

ing to the diffraction pattern. Fig. 3 shows the

diffraction pattern with (yiobs) and (yi

cal) correspond-

ing to two-stage tempered condition. Further details

of this applied technique can be found in Ref. [15].

According to the Rietveld analysis, the double-

tempered PWHT condition contained the higher

amount of retained austenite content (29.2 vol.%).As can be observed in Table 4, the higher austenite

contents are associated with the higher values of

impact toughness.

Carbonitrides of the type M2(C,N) were also

identified by means of X-ray diffraction in both single-

and two-stage tempered conditions. Fig. 4 shows the

diffraction pattern of the two-stage tempered condi-

tion, where the peaks corresponding to carbonitrides

have been noted. The presence of carbonitrides instead

of carbides of the type M7C3 or M23C6 is associated

with the high nitrogen content of the weld metalstogether with the presence of molybdenum. Usually

during tempering, the carbon is precipitated as car-

bide, and according to Irvine [16], precipitates such as

M3C, M2(C,N), M7C3, and M23C6 appear. According

to Pickering [17], the presence of nitrogen (and

molybdenum) promotes the formation of M2(C,N) at

the expense of M7C3 or M23C6.

Thus, both single- and two-stage tempering lead

to martensite decomposition together with the pre-

cipitation of a very thin austenite dispersion, since it

is known that the precipitation of austenite takes

Fig. 5. SEM micrograph of retained austenite between the

martensite laths. Single-stage tempering; 3000, magnifi-cation bar is 10 mm.

Fig. 6. SEM micrograph of retained austenite between the

martensite laths. Two-stage tempering; 3000, magnifica-

tion bar is 10 mm.

Fig. 7. Line scan analysis of carbon on an austenite particle;

20,000, magnification bar is 0.2 mm.

Table 5

Average chemical composition of austenite particles as

measured by TEM/EDS (wt.%)

Cr Ni Mn Si Mo Fe

13.7 8.3 1.9 0.8 1.3 balanced



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place at tempering temperatures slightly higher than

Ac1 (600C). This austenite is shown in Figs. 5 and 6,

for single- and two-stage tempered conditions respec-

tively, arrayed like platelets among martensite laths. It

remains stable and does not later transform intomartensite during the cooling after the tempering.

The line scan analysis of carbon on an austenite

particle in the single-stage tempered condition, deter-

mined by EPMA, is shown in Fig. 7. This indicates the

presence of a considerable amount of carbon. The

actual content level has been determined by X-

ray diffraction from lattice parameter measure-

ments. The austenite lattice parameter measured was

aog= 0.35901 nm, from which the carbon content has

been calculated using Eq. (1) from Refs. [18,19]:

aog 3:572 0:033 wt:% C: 1

According to this equation, the carbon concentration

is calculated to be 0.548 wt.%.

The approximate composition of the austenite

phase as evaluated by EDS analysis using TEM is

presented in Table 5. This indicates that the austenite

particles are enriched in solute elements such as Ni,

Mn, and Mo, when compared with the nominal

composition of the welding wire (Table 1). This

solute enrichment in the austenite particles will leadto a decrease in the Ms temperature. The Ms tempera-

ture at which the precipitated austenite would begin

to transform spontaneously into martensite has been

calculated in terms of the austenite composition by

Eq. (2), which is specific to soft martensitic stainless

steels (from Ref. [7]).

MsC 492 125 wt:% C 65:5

wt:% Mn 10 wt:% Cr

29 wt:% Ni 2

The Ms temperature thus calculated is 78C.

The microstructure after single-stage tempering

observed by TEM is shown in Figs. 8 and 9. The

fine austenitic platelets are the darkest platelet-like

regions with an increased dislocation density. Other

investigators have observed a similar substructure

with localized carbide particles and orientation rela-

tionships between matrix and austenite, such as the

KurdjumovSachs relationship, in similar soft mar-

tensitic stainless steels [20,21]. All these character-istics suggest that the austenite could have been

formed by a shear process [22].

The Mossbauer spectra of a specimen cooled at

different temperatures down to 196C and held atthis temperature for 20 h are shown in Fig. 10. All

spectra reveal six lines belonging to the different

ferromagnetic phases and a central line, corresponding

to a paramagnetic phase. This central peak corresponds

to the austenite phase. No changes in the austenite

content were noted with decreasing temperature. This

indicates that the austenite particles did not transforminto martensite and retained their thermal stability even

after cooling to 196C.The Ms determined for the enriched retained

austenite was 78C. This value is not in agreementwith the results obtained by subzero treatments.

Although other equations were used to calculate Ms,

the results were similar to that obtained using Eq. (1).

This indicates that the solute enrichment of the

austenite particles could be a factor contributing to

the stability of the particles, but it is not the only

factor accounting for the austenite thermal stability.Accordingly, it is suggested that the stability of the

austenite has substructural as well as chemical ori-

gins. Therefore, there is a possibility that the

observed substructure in the austenite could increase

its stability against the transformation into martensite

on cooling. The stability of the austenite particles is

probably due to increased difficulty in propagatingFig. 8. TEM micrograph of retained austenite between the

martensite laths; 43,000, magnification bar is 1 mm.

Fig. 9. High dislocation density within the retained

austenite particles; 75000, magnification bar is 0.5 mm.



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the shear of the martensitic transformation via the

increased barriers afforded by the substructure in the

austenite particles.

In order to test the mechanical stability of the

austenite particles, some part of the material was

cold rolled at room temperature to reductions inthickness of 20%, 40%, 60%, and 80%. The

mechanical stability of the austenite is defined as

the susceptibility of austenite to transform into

martensite, as induced by plasticity. An austenite of

higher mechanical stability needs more strain to

transform into martensite than that of lower mechan-

ical stability. The changes in the austenite content

associated with the austenite (fcc) to martensite (bcc)

transformation were assessed by Mossbauer spectro-

scopy and are shown in Fig. 11. The spectra show

that with only 20% reduction in thickness some ofthe austenite has been transformed into martensite.

Above 60% reduction in thickness the central lines

corresponding to austenite are absent, so by this

stage practically all the austenite has been trans-

formed mechanically into martensite.

After Charpy V-notch impact testing, regions

close to and far from the fracture surfaces of the

Charpy V-notch specimens in the single-stage tem-

pered condition, were analyzed by X-ray diffraction

to study the interaction between the fracture front

and the austenite particles. This interaction could

result in the transformation of the austenite particlesinto martensite. The X-ray diffraction patterns of

these regions corresponding to single-stage tempered

condition are shown in Fig. 12a and b. The austenite

diffraction peaks were observed in regions far from

the fracture but not in regions close to the fracture

surface. This indicates that austenite particles were

transformed into martensite in regions close to the

fracture. In order to view effectively the microstruc-

ture immediately adjacent to the fracture, the fracture

surface of the Charpy specimen was electrolytically

covered with metallic chromium. An intense strainfield was observed close to the fracture by SEM

(Fig. 13) together with microvoids around trans-

formed austenite particles (Fig. 14). The average

microhardness in this region adjacent to the fracture

Fig. 10. Mossbauer spectra before and after the subzero

treatment.

Fig. 11. Mossbauer spectra before and after cold rolling at

room temperature.



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Fig. 12. X-ray diffraction patterns performed close to and far from the fracture surfaces of the Charpy-V notch specimens for

single-stage tempering: (a) Region far from the fracture surface of Charpy-V notch specimen, (b) Region immediately close to

the fracture surface of Charpy-V notch specimen.

Fig. 13. Intense strain field closer to the fracture surface of

Charpy-V notch specimen; 3000, magnification bar is10 mm.

Fig. 14. Microvoids around transformed austenite particles

close to the fracture of a Charpy specimen; 15,000,

magnification bar is 1 mm.



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surface was 320 HV, which is near the value of the

as-welded condition where the microstructure ismainly martensite. However, far from the fracture,

the average microhardness was 275 HV, in accor-

dance with the value of the single-stage tempered

condition (Table 4).

All these results suggest that the austenite retained

particles suffered the martensitic transformation dur-

ing the progression of a crack front and they could

have acted as energy absorbers. In the event of a

propagating crack passing into or near metastable

regions, the concentrated strain field at the crack tip

enables austenite particles to transform into stable,but less dense, martensite. This transformation,

mechanically induced in the plastic zone, absorbs

additional energy, thus effectively enhancing the

toughness. The associated volumetric expansion of

this transformation tends to close the crack and

relieve stresses at its tip, absorbing strain energy

during the fracture that might otherwise have gone

through the crack extension. This kind of transforma-

tion mechanism is recognized to be primarily respon-

sible for the beneficial toughening effect of a

metastable phase within the microstructure [23].Fig. 6 shows the microstructure of the material in

the double tempered condition (first at 670C/2 h/air,second at 600C/8 h/air). A structural refinement can

be noted, meaning a higher austenite content together

with a more uniform distribution of this phase. The

mechanism by which austenite precipitation increases

with double tempering is thought to be associated

both with the instability of the austenite particles

during cooling following tempering at 670C and tolonger times of tempering at 600C. The austenite thatwas produced at 670C was not stable enough so it

became partially transformed into martensite during

cooling in air. Fig. 15 shows the diffraction pattern of

a X-ray diffraction analysis performed on a specimen

after the first tempering at 670C. The principal peak

of austenite is barely apparent. In this condition the

microstructure shows a second precipitated phase

with similar morphology like the austenite particles

(Fig. 16). In fact, this precipitated phase is mainly

Fig. 15. Diffraction patterns of a X ray diffraction analysis performed on a specimen after the first tempering at 670 C/2h/air.

Peaks of austenite barely appear.

Fig. 16. Precipitated martensite with similar morphology to

the austenite particles, resulting from the martensitic

transformation of the austenite particles which are not stable

enough, after the first tempering at 670C/2h/air; 3000,

magnification bar is 10 mm.



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martensite, resulting from the martensitic transforma-

tion of the only partially stable austenite particles.

Thus, after the first tempering, the microstructure was

composed of tempered martensite, some retained

austenite, and some fresh martensite. During thesecond tempering at 600C, new stable austenite

particles platelets precipitate through new interfaces

(fresh martensite/austenite) and the fresh martensite is

decomposed into tempered martensite. A similar

mechanism has been proposed by other investigators

in these steels [20]. On the other hand, longer temper-

ing times, 8 h, at 600C (four times longer than in

single-stage tempering) promoted a very high content

of retained austenite, more than in single-stage tem-

pered condition. Hence, as result of this two-stagetempering, an increase in the number of austenite

platelets after double tempering was produced (29.2

vol.% according to Table 4) and with a more uniform

distribution. A scheme of this structural refinement

due to two-stage tempering is shown in Fig. 17.

4. Conclusions

1. The studies have shown that the microstructure

of 13% Cr NiMo weld metals, after single- andtwo-stage tempering, consists of tempered martensite

and retained austenite with an acicular or lath-like

morphology, very thinly spread along the martensite

laths. These austenite particles are enriched in C, Ni,

and Mn, and have a high dislocation density. The

lath-like austenite morphology and the substructure

of the austenite phase, suggest that the austenite

par ticles could be for med by a shear process

between the laths of martensite that have been

enriched in stabilizing elements during the intercri-

tical tempering.2. According to the results from treatments at

different subzero temperatures, the austenite particles

were thermally stable. However, they could be

transformed mechanically into martensite by cold

rolling. Mechanical transformation was also observed

close to the crack front of the Charpy V-notch

specimens where localized plastic deformation took

place. This indicated the austenite had been trans-

formed by TRIP.

3. The solute enrichment of the austenite parti-

cles may be a factor contributing to the low Ms ofthe particles, but it does not fully account for the

high austenite thermal stability. The presence of a

high dislocation density within the austenite parti-

cles suggests that the stability of these particles may

have substructural as well as chemical origins.

Some stability of the austenite particles may be

probably due to increased difficulty in propagating

the shear of the martensitic transformation via

increased barriers afforded by the substructure in

the austenite particles.

4. The yield and tensile strength values werehighest for the as-welded condition. By comparison,

the greatest softening was induced by the PWHT that

involved a two-stage temper treatment. This heat

treatment also produced the largest amount of

retained austenite. Regarding the mechanical proper-

ties, a good balance was achieved between maximum

toughness, high ductility, and high strength.

Fig. 17. Scheme of the structural refinement due to two-stage

tempering: (a) 950C/1h/air, Martensite lath (MI), (b) 950C/

1h/air + 670C/during heating. Tm: tempered martensite

g: austenite, (c) 950C/1h/air + 670C/2h/air (after cooling).

Tm: tempered martensite; MI: Martensite lath, g: austenite,

(d) 950C/1h/air + 670C/2h/air + 600C/2h/air. Tm:

tempered martensite; g: austenite.



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characteristics and effects of austenite resulting from

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