In-situ Investigation during Tempering of a High Speed Steel with X-ray Diffraction

$Page 1: In-situ Investigation during Tempering of a High Speed Steel with X-ray Diffraction$
In-situ Investigation during Tempering of a High Speed Steelwith X-ray Diffraction

Manfred Wießner*, Siegfried Kleber**, Alfred Kulmburg***(Received: 17 January 2006)

DOI: 10.1002/ppsc.200501002

1 Introduction

High speed steels are tooling steels, which are mainlyused for cutting processes and conserve their extremehardness also during red heat. They obtain their optimalproperties by hardening at 1200 °C followed by repeatedtempering between 500 °C and 600 °C. After hardeningthe microstructure contains eutectic carbides, martensiteand retained austenite which varies from 25 to 35 %.While retaining at tempering temperature, carbides aredeposited from the residual austenite and are in therange of 10 nanometers. These carbides cause precipita-tion hardening. Alloying additions decrease, and as aresult the residual austenite gets instable and is trans-

formed to martensite. The so-called secondary harden-ing causes a hardening maximum at the temperaturerange between 540 °C and 570 °C, depending on thetempering temperature and time. The hardening maxi-mum is very stable and does not decrease even afterrepeated tempering at this range explaining also the redhardness.

2 Experimental Set-up

The processes during secondary hardening are investi-gated in view of the composition of the carbides (c.f. stu-dies from Karagöz and Andrén [38]) and dilatometri-cally in the context of phase transformations (c.f. studiesfrom Kulmburg et. al. [39,40]). The used methods werenot suitable to estimate changes in lattice parameters ofmartensite and austenite. Therefore, X-ray diffractionwas used to observe the processes during tempering in-situ.Conventionally cast material tends to segregate duringsolidification causing anistropic distribution of carbidesand segregation in the matrix. In order to minimizethese effects, a powder-metallurgical (PM) manufac-tured high speed steel Boehler S290PM was selected.The microstructure of this PM-steel consists of eutectic

Part. Part. Syst. Charact. 22 (2005) 407–417 407

* DI M. Wießner, Böhler-Edelstahl GmbH, Mariazellerstraße25, 8605-Kapfenberg (Austria).E-mail: [email protected] (correspondingauthor)

** Dr. DI S. Kleber, Böhler-Edelstahl GmbH, Mariazeller-straße 25, 8605-Kapfenberg (Austria).E-mail: [email protected]

*** Univ. Doz. Dr. mont. A. Kulmburg, TU Graz, Institut fürWerkstoffkunde, Schweißtechnik und Spanlose Formgebungs-verfahren, Kopernikusgasse 24/I, 8010-Graz (Austria).

Abstract

The combined use of a X-ray diffractometer and a hightemperature chamber allows in-situ determination ofmicrostructure during heat treatment. X-ray diffracto-grams are recorded during tempering of high speedsteels. The Rietveld method is used in connection withsize-strain analysis. Lattice parameters of austenite andmartensite are utilized for estimating the change of car-

bon content that has massive influence on the secondaryhardening. Nevertheless, it was shown that the heattreatment of deep-frozen high speed steel cannot bededuced a priori from the applied tempering processunder conventional quenching conditions. The differ-ently conditioned microstructures are based on differentphase amounts and microstrain-situations.

Keywords: Carbon content, High speed steel, Size-strain analysis, Tempering, X-ray diffraction

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408 Part. Part. Syst. Charact. 22 (2005) 407–417

carbides with a diameter between 3 and 5 lm, isotro-picly embedded in a homogeneous composed matrix.This material does show almost no heterogeneity.Tempering heat treatment was carried out in a preciselycontrollable high temperature chamber (HTK). Duringtempering the X-ray diffractograms are recorded inequidistant time steps. Analysis of the diffractogramswas carried out in consideration of the influence of themeasurement system. The integral breadth of the dif-fraction reflexes was used for implementing „size-strain“-analysis which allows for characteristic para-meters despite the overlapping of diffraction lines frommartensite, austenite and carbides. This tempering in-vestigation was performed on a water-quenched sampleat room temperature and a specimen which wasquenched in fluid nitrogen. The parameters that weredetermined through X-ray diffraction were interpretedwith models from the technical literature.

3 Specification of the Measurement Device

The X-ray measurements were made by a D8 Advancediffractometer from the Bruker AXS company. A con-figuration was used which differs from a classicalBragg-Brentano-device. On the primary side of the D8 aGoebel mirror generates a parallel beam. If the sampleis not aligned, which cannot prevented during temper-ing, also high intensities will be reached which is anadvantage in comparison to Bragg-Brentano configura-tions [1].For the investigation a chrome tube was used. The wave-length of Chrome-Ka-radiation is by far longer as forexample of copper-Ka-radiation, which makes a separa-tion from overlapped peaks much easier. Furthermore,when copper-Ka-radiation is used, iron shows strongX-ray fluorescence [2] and permits only a small penetra-tion so that possible marginal oxidation has a negativeeffect during heat treatment. Detection of the X-rayswas carried out with a Position Sensitive Detector whichsimultaneously detects X-rays over a angular range ofD2h = 8°. The gain of time when compared to a szintilla-tor is acquired by line broadening and asymmetries ofthe diffraction peaks [3]. The high temperature systemHTK 1200 from the company Anton Paar was used asthe heating system [4]. This chamber is transparent forX-rays along a special “window”. As a result of thermicexpansion of the sample holder the position of the speci-men is changed during the experiment [5]. Nevertheless,the parabolic curved Goebel mirror makes it possible toreach high diffraction intensities.

4 Determining the Influence of the Diffractometer

Lattice defects lead to non-exact fulfilment of destruc-tive interference which causes broadening of the diffrac-tion peaks [6]. But a fraction of the peak breadth can beassociated with the instrumental system. The Rietveldmethod was used which is suitable for estimating quanti-tative phase amounts, lattice parameters and lattice de-fects for the measured X-ray diffractograms. To get sui-table precise values knowledge about peak broadeningin consequence of the used instrument is necessary. Themodeled diffractogram Y is expressed as a convolutionof instrument profile G and a broadening model of thesample S (see Eq. (1) – a symbolic representation). Thetotal instrument profile G can be described as a functionof the emission profile W of the tube and a geometricalaberration B [7–9].

Y 2h� � � G × S � W × B� � × S �1�

The influence of the diffractometer can be estimatedwith a “nearly” ideal sample, which has hardly latticedefects and shows furthermore a satisfactory number ofdiffraction peaks. Mostly at small diffraction angles the“geometrical aberration” influences the instrument pro-file, and at wide angles the emission profile exerts influ-ence [8]. This effect is estimated by the NIST 660a Stan-dard (Lanthanum Hexaboride powder). The latticeparameter is known, therefore it is possible to simulta-neously calibrate line form and line position [37]. Thispowder also has sufficiently high number of diffractionpeaks for chrome-Ka-radiation.Robert W. Cheary and Alan A. Coelho developed a fun-damental model for Bragg-Brentano-diffractometers,which describes peak broadening as a function of thegeometrical settings [7,10,11]. For these investigations aBragg-Brentano-model is required which causes thesame peak broadening compared to the used measure-ment system. The Goebel mirror can change the wave-length composition of the chrome tube. By this the exactwavelength distribution was also refined.The software package Topas 2.1 uses this fundamentalparameter approach and is therefore capable of estimat-ing the instrumental influence. The parameters of thismodel are varied for as long as the difference betweenthe LaB6-measurements and the modeled diffracto-grams gets a minimum (see Figure 1).The used fundamental parameter approach for Bragg-Brentano diffractometers contains a suitable amount ofmodel parameters and functions and can also be usedfor diffractometers that differ from the classical Bragg-

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Brentano configuration. However, the fit parameterslose their physical meaning.

5 Line broadening as a Consequence of LatticeDefects

Generally, each perturbance of the ideal crystal latticeleads to line broadening. Division of a crystal into small,incoherent, diffracting domains (small-angle boundaries,fault stacks, twins or other extended imperfections),leads to line broadening as a result of the finite domainsize. Dislocations, vacancies, interstitials or substitu-tionals lead to strain broadening [12]. Generally, bothphenomena are coupled, complicating physical interpre-tation.Already Scherrer discovered the inverse correlationbetween line broadening and domain size. Stokes andWilson conducted general investigations of the peakbroadening ßc (full width at half maximum) resultingfrom finite domain size distribution, which is indepen-dent of crystal form and symmetry (Eq. (2)).

bc � k��Lvol � cos h� �2�

Dislocations, vacancies, interstitial and substitutionalscause strain broadening (see Figure 2) as mentioned

above. The maximum lattice strain e (Eq. (3)) is deducedfrom the volume weighted width Dd of the lattice spacingd [13].

e � Ddd

�3�


Fig. 1: Comparison between measuredand fitted diffractogram (with use of thefundamental parameter approach – 001and 311-reflex)

Fig. 2: Example for lattice strains (analogous [8,14]).


The peak broadening ße (full width at half maximum)caused by strain broadening is displayed in Eq. (4).

be � 4 � e � tan h �4�In Topas 2.1 the Double-Voigt approach is implemen-ted, which combines peak broadening caused by finitedomain size and strain broadening. The peak profilesare described as convolutions of Gauss and Lorentzfunctions to explain domain size Lvol and maximum lat-tice strain e. The full width at half maximum depends on1/cos(h) for the size-part and tan(h) for the strain-part[6,8,15,16].For martensite and austenite only a small amount of dif-fraction peaks are available. Therefore the reflexes areenlarged due to the used chrome-Ka-radiation com-pared to copper-Ka-radiation. But this effect facilitatesto separate the diffraction peaks from austenite, marten-site and eutectic carbides. The small amount of reflexesleads to more fluctuant microstructure parameters dueto the increased correlation between the parameters ofthe Double-Voigt model. Furthermore, dislocations areproduced during the phase transformation from auste-nite to martensite [17]. But those are inadequatelydescribed by this model [18].

6 Implementation and Measurement Results

As test material cylindrical samples with 16.5 mm dia-meter and 2.4 mm height are used. The specimens of thesteel grade Boehler S290PM have the following chemi-cal composition (in weight percent):

The original material is available in an annealed con-dition. The austenitizing temperature is 1200 °C and thesoak is chosen at 15 minutes.Then a sample is water-quenchedto room temperature and asecond one in liquid nitrogen(–196 °C).The microstructure consists of fer-rite and different carbides fromthe type M23C6, M6C and M8C7 inannealed condition and of mar-tensite, (retained-) austenite andcarbides from the type M6C andM8C7 in quenched condition. Thediffractograms are characterizedby peak overlapping for the mar-tensite, austenite and carbide

reflexes. The time temperature curve is presented inFigure 3 and the measurements are made during iso-thermal steps. In order to get reasonable and analysablediffractograms, measurement time was chosen to 3512seconds. Heating and cooling between isothermal stepsoccured linearly.Figure 4 shows an in-situ diffractogram analysed withthe software package Topas 2.1 for the water-quenchedsample tempered at 100 °C.

6.1 Martensite

The typical splitting of the martensite peaks is caused byinterstitially dissolved carbon. Simultaneously, an enor-mous amount of lattice defects generates broadening ofthe diffraction peaks resulting in an overlap of the pairwise reflexes (Figure 5) [19].The estimated tetragonality of the martensitic phase is ameasure for the interstitial dissolved carbon content perunit cell [20,21]. Furthermore, a high dislocation densityis built during austenite to martensite transformation.The high dislocation density of martensite shows a


C Si Mn P S Cr Mo V W Co2,00 0,43 0,30 0,021 0,018 3,77 2,49 4,83 14,27 11,00

Fig. 3: Time temperature curve in the high temperature chamberHTK 1200.

Fig. 4: Analyzed diffractogram by Topas 2.1, water-quenched, tempering 100 °C.


strong peak-depended anisotropic full width at halfmaximum [18,22]. The investigations were carried outwithout under consideration of this effect and only anisotropic average lattice strain e was used. Since the fullwidth at half maximum of the diffraction peaks can dif-fer from Eq. (2) and Eq. (4), the modeling of the mar-tensite peaks is eventually afflicted with larger errors.The tetragonal deformation of the unit cell correlateswith the dissolved carbon content. This is an importantparameter since the dissolved carbon content is neces-sary for formation of secondary hardening carbides. Butthese nm small carbides could not be directly detectedin the diffractograms. In Figure 6 the lattice parametersof martensite are displayed for heating up to 580 °C andthe following cooling down as a function of the tempera-ture. The lattice parameters a and c of martensite wereestimated from the in-situ measured diffractograms.It is striking that the deep-frozen specimen in relation tothe sample which was quenched to room temperatureshows a larger tetragonality (see Figure 6). The para-meters acquired by X-ray diffraction reflect an averagevalue for the complete martensitic phase. The tetragon-ality of individual martensitic domains depends on thetransformation temperature because the transformationrate is a function of the induced strain and the local car-bon content. At lower temperatures transformed mar-

tensitic domains shows a higher dissolved carbon con-tent which affects also the local tetragonality.Furthermore, the cooling rate for the deep-frozen speci-men was much higher than for the specimen quenchedin water one so that auto-tempering effects whichreduce also the tetragonality during cooling down, arelower for the deep-frozen sample.During tempering tetragonality starts to decrease at atempering-temperature of 100 °C. In the literature thiseffect is ascribed to the precipitation of e-carbides[23,24].The carbon atoms are located in positions which are en-ergetically disadvantageous. But the low diffusion velo-city allows only accumulation of coherent clusters. Inter-estingly the tetragonality of the deep-frozen sampleremains at a higher level for the whole heating-up pro-cess, in contrary to the water-quenched sample. Thedeep cooling possibly induces more lattice imperfectionsduring micro ductile processes due to the higher internalstrains. In any case, the deep-frozen sample showshigher maximum lattice strain e during heating up ascompared to the water-quenched sample (see Figure 7).Measurable peak broadening in consequence of thefinite domain-size was not detected (as in [17]) which isin accord to micrographs (for example also electronbackscattering diffraction). Possible high misorientations


Fig. 5: Splitting of the martensite peaks.

Fig. 6: Lattice parameters of martensite during temper-ing for the water-quenched and the sample quenched inliquid nitrogen.


in the martensite laths are caused by a high amount ofdislocations, which affects the increase in the solely iso-tropic estimated maximum lattice strain e.Due to the fact that the lattice parameters are superim-posed by thermal expansion it makes sense to define arelative tetragonality rt (Eq. (5)) to get a measure of thedissolved carbon content. The relative tetragonality rt isdirectly proportional to the interstitial dissolved carboncontent [20,21].

rt � c � ac

�5�

Figure 8 shows for both specimens the relative tetragon-ality rt for heating-up and cooling down as a function ofthe temperature during tempering.Between 200 °C and 500 °C for both specimens the tetra-gonality is temperature independent within the accuracyof measurement.In the range of secondary hardening above 500 °C thetetragonality decreases due to the formation of morenm scaled carbide precipitations. At 580 °C both sam-ples show the same lattice parameters (Figure 6 and 8).During cooling there are only small measurable differ-ences between both quenched conditions. The water-quenched sample builds more martensite due to thehigher amount of residual austenite during cooling. Thedifference between the two quenched samples is ap-

proximately 10 w.-%. Also the maximum lattice strainfor both heat treatments is similar during cooling down.During this process, in this phase the water-quenchedsample shows a slightly rising lattice strain for tempera-tures below 300 °C (see Figure 7). In the next section itis shown that this is in agreement with the residual aus-tenite transformation. Therefore the newly built marten-site is accountable for these phenomena because of thehigher lattice defect than in tempered martensite. InFigure 9 the austenite amounts of both samples areshown as a function of temperature during heating-upand cooling down.

6.2 Residual Austenite

The effects during tempering between the tetragonaldistorted martensite correlate with phenomena of thecubic-face centered austenite phase. For this reasonphase amounts (see Figure 9) and lattice parameter a(see Figure 10) of the residual austenite are also esti-mated.It is remarkable that the increase of the lattice para-meter is disproportionate to thermic dilatation. In differ-ent literature sources is presented that the residual aus-tenite additionally dissolves interstitial carbon duringheating [24,25,27–31]. Deviant effects are not discussed.Figure 11 shows the carbon distribution at a tempered


Fig. 7: Maximum lattice strain of martensite duringtempering for the water-quenched and the samplequenched in liquid nitrogen.

Fig. 8: Relative tetragonality of the martensite duringtempering for the water-quenched and the samplequenched in liquid nitrogen.


martensite-austenite phase interface for a model inwhich part of the carbon dissolved in martensite is“pushed” into the residual austenite.As a consequence in the region of an interface a locallyvery high concentration exists, leading to the formationof e-carbides. These carbides are energetically unba-lanced, therefore by further tempering they can neitherbe resolved nor converted to more stable modifications[23,24]. The increased carbon concentration at the phaseinterface causes in any case an effective diffusion flowof carbon into the residual austenite.

Carbon has a higher diffusion coefficient in the tetrago-nal body-centred martensite than in the cubic face-cen-tered austenite. If the carbon concentration is nearly thesame in both phases it is more likely that a carbon atomdiffuses into the austenite than vice versa. Such phe-nomena are also observed during the formation ofbainite [32,33]. Due to the fact that the amount of resi-dual austenite is lower for the deep-frozen sample, theaustenite domains are smaller and the ratio betweenphase interface to phase volume is much higher. For thissample more carbon per volume diffuses in the austenitethan for the water-quenched sample. Additionally, thedriving force of diffusion is much higher for the samplequenched in liquid nitrogen so that the disproportionateincrease of the lattice parameter starts at lower tem-peratures.But the water-quenched sample reaches this level notbelow 200 °C. The lattice parameter of residual austenitedecreases at 400 °C for the deep-cooled sample and at450 °C for the water-quenched sample, respectively. Thisevidence is deduced from the findings that also in resi-dual austenite carbides are precipitated [34].In [35] the relation between the change of lattice para-meter, temperature, and dissolved amount of carbon isdetermined. If this model is used for the S290PM thefollowing relative carbon trend can be displayed (see


Fig. 9: Amount of residual austenite during temperingfor the water-quenched and the sample quenched inliquid nitrogen.

Fig. 10: Lattice parameter of residual austenite duringtempering for the water-quenched and the samplequenched in liquid nitrogen.

Fig. 11: Distribution of carbon at a tempered martensite-austenitephase interface (figure similar [26,32], see also [33]).


Figure 12). The origin of the scale is arbitrarily chosenwith 0 for the water-quenched sample.During tempering heat treatment the residual austenitedecreases massively above 300° during cooling downand it is transformed to martensite. After tempering theretained austenite is below the measurable detectionlimit.

6.3 Carbides

The material characteristics for applications are onlymet with a considerable phase amount of eutectic car-

bides [34,36]. When tempering a nearly linear increaseof the lattice parameter can be observed for the fractionof the eutectic carbides, which is a consequence of ther-mal dilatation (see Figure 13 and 15).Furthermore already during heating-up a decrease ofthe lattice microstrain is measurable indicating the heal-ing of lattice defects (see Figure 14).After cooling it is noticeable that the original latticeparameter will not be reached. Maybe this effect is incorrelation with the reduction of microstrain but also asmall change of the stoichiometrical composition cannotbe excluded. In the literature some authors use a slightlymodified structural model for the M8C7 carbide, forexample with the half lattice parameter for the unit cell.


Fig. 12: Relative carbon content of the residual auste-nite (relating to the water-quenched sample at roomtemperature) during tempering for the water-quenchedand the sample quenched in liquid nitrogen.

Fig. 13: Lattice parameter of M8C7 during temperingfor the water-quenched and the sample quenched inliquid nitrogen.

Fig. 14: Maximum lattice strain of M8C7 during temper-ing for the water-quenched and the sample quenched inliquid nitrogen.


A similar behaviour of the lattice parameter (seeFigure 15) and the microstrain (see Figure 16) in con-nection with tempering heat treatment was observed forthe M6C-carbide.

7 Summary

The combined use of a X-ray diffractometer and a hightemperature chamber allows – in contrast to a dilato-metric method – in consequence of a change of micro-structure during heat treatment to directly attribute theseparate phases.The used analysing algorithm is suitable for complexcompositions of phases. Nevertheless, it was shown thatthe heat treatment of deep-frozen high speed steel can-not be deduced a priori from the applied temperingprocess under conventional quenching conditions. Thedifferently conditioned microstructures are based ondifferent phase amounts and different microstrain-situa-tions can be assumed.

8 Nomenclature

h degrees diffraction angleW emission profileB geometrical aberrationHTK high temperature chamberG instrument profilea 10–10 m lattice parameterc 10–10 m lattice parameterd 10–10 m lattice spacinge – maximum lattice strainY modeled diffractogrambc rad peak broadening finite domain sizebe rad peak broadening strainPM powder-metallurgicalrt – relative tetragonalityS sampleDd m volume weighted lattice spacingk m wavelengthLvol m domain size (volume weighted average)


Fig. 15: Lattice parameter of M6C carbide during tem-pering for the water-quenched and the samplequenched in liquid nitrogen.

Fig. 16: Maximum lattice strain of M6C during temper-ing for the water-quenched and the sample quenchedin liquid nitrogen.


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In-situ Investigation during Tempering of a High Speed Steel with X-ray Diffraction