études et recherche - · PDF fileand deposit cracking susceptibility. Standard...

37 I janvier-février 2017 I Soudage et techniqueS connexeS

études et recherche

Fabrice SCANDELLA1 and

(INSTITUT DE SOUDURE GROUPE)

Jean-Marie BONNEL2

(WELDING ALLOYS FRANCE)

This paper summarises theoretical and experimental work undertaken to develop a hardfacing product designated “167Nb-S”. This material, which is based on the Fe-Cr-Nb-C system, was designed for cladding components subjected to severe abrasive wear. The work undertaken includes microstructural characterisation, solidification modelling and abrasion testing. The microstructure of 167Nb-S consists of a large volume fraction of primary niobium carbides, surrounded by a martensitic matrix containing fine quasi-eutectic niobium carbides. Some retained austenite is found along grain boundaries, together with small amounts of intermetallic phase. Alloy development enabled a reduction in the amount of retained austenite and intermetallic phase present in early versions of the alloy, which is beneficial in terms of abrasive wear resistance and deposit cracking susceptibility. Standard three-body abrasion tests showed that 167Nb-S exhibits wear rates comparable to those of high carbon/high chromium hardfacing deposits with similar hardness. 167Nb-S offers many advantages over conventional hardfacing products as high quality multi-layer deposits with hardness up to 61 HRC can be deposited crack-free, and these remain machinable. A highly successful application of this material is the cladding of roller press rollers, as an alternative to titanium-containing hardfacing consumables or segments.

1. List of symboLs

d{hkl} Spacing between planes with indices h, k and l

r Average wear ratet Timew Wear parameterxδ Mole fraction of δ ferriteD Rubber-coated wheel diameterH Hardness (Rockwell C scale)HV Vickers hardness numberL LiquidM Mass of clinker groundNbC(p) Primary niobium carbideNbC(q–e) Quasi-eutectic niobium carbideα Ferriteα' Martensiteγ Austeniteγret Retained austeniteδ High temperature body-centred cubic form

of ferriteθ Celsius temperatureρ Density of hardfacing materialsΔm Mass lossΔV Volume lossΔVadj Adjusted volume loss Θ Angle between the specimen surface and

the X-ray source or counter

Development of a harDfacing material in the fe-cr-nb-c system for use unDer highly abrasive conDitions

1. Institut de Soudure, [email protected]. Welding Alloys France, [email protected]

2. introduction

2.1 WEAr rESiStANt hArDFACiNg prODuCtS

Material and energy conservation is becoming increasingly important. Besides fatigue and corro-sion, wear has been identified as a major cause of material wastage; any reduction of wear can the-refore result in considerable savings. In situations where improved wear resistance is required in com-ponents to extend their service life, the optimum solution is to apply a wear resistant surface onto a cheaper and tougher core. Among numerous surface treatments available, hardfacing using tubular wires is one of the most versatile processes, enabling a wide range of materials to be deposited onto most ferrous surfaces. The range of commercial hardfa-cing products is extremely large; the materials listed in Table 1 are standard highly alloyed iron-base hardfacing consumables manufactured by Welding Alloys, for use under highly abrasive conditions when the working temperature remains below about 300 °C. Though very different in terms of micros-tructure, the hardfacing products listed in Table 1 share two inherent features: high cracking sensiti-vity and lack of machinability. Stress relief cracking occurs upon solidification of these materials whene-

Product designation

AWM composition, wt%. Balance: FeH, HRC Typical cladding

applicationsC Si Mn Cr Mo V Others

HC-O 5.0 1.5 1.5 27 … … … 61 Mixers, wear plates

TiC-O 1.8 0.7 1.2 6.5 0.8 0.2 5.0 Ti 57 Hard material crushers

BN-O 0.5 1.3 2.0 … 2.0 … 4.5 B 65 Earth-moving equipment

CN-O 5.0 1.0 0.5 22 … … 7.0 Nb 63 Riddle plates

VN-O 5.0 1.2 0.5 23 … 10 … 63 Mixer paddles, crusher armouring

CARBW-O 50 to 60 wt.% WC 63 Dust extraction ventilators in the mining and cement industry

Table 1: Three-layer deposit chemical composition, hardness and typical applications of some commercial hardfacing products [1]

1612_0347_P_037_048_ETUDES_ET_RECHERCHE.indd 37 26/01/2017 11:15

38 Soudage et techniqueS connexeS I janvier-février 2017 I

CONJONCturEétudes et recherchever deposited metal is even slightly constrained by the substrate. Milling or turning of such hardfacing deposits is impossible as it quickly leads to carbide tip or ceramic tool break-off.

2.2 ABrASivE WEAr

The concise introduction provided here is restricted to the situation of three-body abrasion, which occurs when hard particles are free to roll and slide between two surfaces in relative motion [2, 3]. The wear situa-tion is mainly governed by abrasive particles proper-ties (hardness, shape and size) and those of the mate-rial against which the particles are rubbing (hardness, microstructure) [2, 3]. The hardness of abrasive par-ticles influences wear rates; particles with hardness close to that of a surface being abraded cause much less wear [3]. However, if the particles are signifi-cantly harder than the surface subjected to abrasion (which often is the case), the exact hardness value of the particles is of lesser importance, as wear rates are high [3]. Wear rates are also strongly dependant on particle shape, as angular particles cause greater wear than round ones [3]. Angular particles have a particularly detrimental effect on carbide-containing alloys due to preferential wear of the matrix and extensive carbide cracking [4]. Many studies have confirmed that wear rates decrease with particle size (a marked drop in wear rates is observed when par-ticles are smaller than about 100 µm) [3, 4].In the case of steel matrices, optimum resistance to three-body abrasion is obtained with high carbon martensite [2, 5]. The presence of additional hard phases such as carbides or nitrides can dramatically improve wear resistance, which is why hardfacing products are highly alloyed materials. The contri-bution of these hard constituents to abrasive wear

resistance is related to their size, spacing, volume fraction and hardness. Their effectiveness will obviously also depend upon the properties of the abrasive. For instance, fine carbides may be dug out within chips from the surface of a hardfacing, while larger carbides have to be worn away and therefore contribute tremendously to wear resistance [6]. A decrease in carbide spacing is beneficial to abrasive wear resistance as the penetration of abrasive par-ticles is more difficult [7]. The hardness of various pure carbides can be compared in Figure 1.MC-type carbides with a face-centred cubic crystal structure (TiC, VC, ZrC and NbC) are significantly harder than other types of carbides. The hardness of these compounds is however not only related to crystal structure, but will depend on the nature and amount of solute atoms (carbides found in complex hardfacing alloys will invariably exhibit some solu-bility for various elements present) [12]. It has been suggested that vanadium and niobium carbides give the best resistance to abrasive penetration in alloy steel surfaces [2]. The strength of the interface between matrix and additional phases is also of importance and it has been mentioned that a strong interface enhances wear resistance [6, 13]. Relating microstructure and wear properties is however very difficult, often leading to controversial results.

2.3 EvALuAtiON OF ABrASivE WEAr rESiStANCE

Hardness has been identified as a dominant requi-rement for three-body abrasive wear resistance [2]. Although hardness and wear resistance are related, the hardness of a given material should not be used to predict its resistance to abrasion [2, 14]. Hardness should only be used as an indication of the wear

resistance of materials within a given alloy system, as long as no important microstructural changes take place. Hard deposits are often rated by means of a wear coefficient, determined by laboratory tests, but even these results may not be representative of the behaviour of a material in service. In-situ experimen-tation is therefore recommended; if impracticable, wear tests should be conducted under conditions close to those encountered in service. Abrasion tests are essential as these enable a quantitative ranking of materials and to some extent, anticipate the resis-tance to abrasion of a component during service.

2.4 rOLLEr prESS rOLLErS

To date, the major application of 167Nb-S is the cladding of roller press rollers, which are used in the cement industry to crush clinker (the major consti-tuents of clinker are CaO and SiO2, roughly in the proportion 3:1). Giving a detailed account of the ope-ration of a roller press is beyond the scope of this paper; this section therefore provides key informa-tion relevant to rollers. A typical roller press layout is shown in Figure 2.Grinding plants adopted roller press technology in order to increase productivity and reduce power consumption [15, 16]. Roller presses have been in use in cement plants for almost 20 years, during which more than 400 presses have been installed [17]. Rollers require hard-wearing surfaces in order to withstand damage during comminution. Roller surface deterioration occurs as a result of two dis-tinct phenomena: (i) gradual surface removal due to three-body abrasive wear with the feed; (ii) sporadic surface breakdown due to so called “tramp metal”. The presence of tramp metal in the feed is known to cause severe roller damage such as crack initiation

Figure 1Average hardness of pure carbides [8, 11].

Figure 2Schematic representation of the crushing operation carried out in a roller press. Note: this particular arrangement consists of a fixed and a mobile roller, the horizontal movement of the latter determining the pressure applied to the material in the roller gap.



and propagation, which often leads to spalling of the lining; magnetic separation is used to reduce such occurrences [17]. The main requirements for roller sur-faces are therefore abrasive wear resistance, strength and durability [17]. Four types of roller linings are used to resist abrasive wear: hardfacing, studs, segments and sleeves [15, 17]. 10 year-old data showed that out of 400 roller presses used in cement plants, about 300 were equipped with hardfaced rollers [17]. For these hardfaced rollers, the ability to perform repairs on worn or spalled regions is of great importance, as is in-situ refurbishment [17].

3. ExpErimEntaL Work

3.1 ALLOy DESigN

Several hardfacing products based on the Fe-Cr-Nb-C alloy system have been designed by Welding Alloys to meet specific requirements; the composi-tional range of existing materials is given in Table 2.

or a weld [19]. Such simulations have previously given very realistic results in the case of the rapid solidifica-tion of high speed steel deposits [20].

3.2 MANuFACturE OF CLADDiNg CONSuMABLES

A highly versatile processing route is used for the manufacture of hardfacing consumables such as 167Nb-S. The final product is a cored wire with dia-meter generally ranging from 1.6 to 3.2 mm. On a basic level, wire making involves mixing of various powders, which are then inserted into a shaped metallic sheath, as illustrated in Figure 3.

(XRD) plots were acquired in the angular range 10° < 2Θ < 140°. Microstructural examination of specimens was carried out using standard light microscopy and scanning electron microscopy (SEM). All specimens were etched using Marble's reagent (4 g CuSO4, 20 cm3 HCl, 20 cm3 distilled water, etching time 10 s). Energy dispersive spectro-metry (EDS) facilities available on an SEM enabled microanalyses of individual phases to be obtained. Image processing and a simple image analysis com-puter program were used to determine the volume fraction of phases. Finally, some abrasion test specimens were selected to examine abraded sur-faces in the SEM, in an attempt to provide a better understanding of wear mechanisms at a microsco-pic level.

3.5 MEChANiCAL tEStiNg

Rockwell hardness tests were performed using a standard diamond indenter (C scale). Vickers micro-hardness tests were carried out under a load of 200 g. Dry sand/rubber wheel abrasion tests were performed according to ASTM standard G65-94 - procedure A [21]. This test, which leads to a quan-titative ranking of materials in their resistance to scratching abrasion under a specified set of condi-tions, was particularly suited for this work. The dry sand/rubber wheel test involves the abrading of a standard test specimen with dry rounded silica sand. The abrasive is introduced between the test speci-men and a rotating wheel with a chlorobutyl rubber rim. The test specimen is pressed against the rota-ting wheel at a specified force by means of a lever arm while a controlled flow of grit abrades the test surface. A photograph of the apparatus is shown in Figure 4, together with a schematic diagram of the set-up. The rotation of the wheel is such that its contact face moves in the direction of the sand flow. Specimens are weighed before and after the test on a balance with a sensitivity of 0.001 g. Mass losses are converted into volume losses using the following equation: ΔV = Δm (1) ρ

Note: Average density of experimental alloys: ρ = 7.675 g cm–3.

The linear abrasion sustained by test specimens however slowly decreases as the diameter of the rubber wheel decreases as a result of wear of the rubber rim. This effect can be accounted for by converting volume losses into adjusted volume losses, as follows:

ΔVadj = ΔV 228.6 (2) D

Note: Initial diameter of a rubber wheel: D = 228.6 mm.

Development of a harDfacing material in the Fe-Cr-Nb-C system for use unDer highly abrasive conDitions

Element C Si Mn Cr Mo Nb Ti V

Minimum 0.6 1.0 0.5 5 … 3 … …

Maximum 1.6 2.5 1.5 10 1.5 10 0.3 0.2

Table 2: Compositional range of alloys based on the Fe-Cr-Nb-C system, wt %. Balance: Fe

The least alloyed material, designated 51Nb-S, contains 0.6 C, 5 Cr and 3 Nb (wt.%). With 1.4 C, 6 Cr and 8 Nb (wt.%), 167Nb-S is significantly more alloyed. The matrices of existing alloys whose com-position is within the range shown in Table 2 are invariably found to be martensitic; variations in the alloying content however have a significant effect upon the nature and volume fraction of carbides, retained austenite and intermetallic phase forming upon solidification. Microstructural differences are briefly discussed in a later section, as these have a fundamental effect upon deposited metal cracking susceptibility and wear properties.Solidification modelling was used not only to inves-tigate the solidification of 167Nb-S, but also to solve problems which arose during alloy development. Given that deposited metal usually experiences high cooling rates, the microstructure obtained (without heat treatment) is far from equilibrium; when applied, heat treatments are not severe enough to enable equilibrium conditions to be reached either. The fact that non-equilibrium solidification occurs therefore had to be taken into account to ensure realistic solidi-fication models were obtained. This led to the choice of the Scheil simulation module of ThermoCalc, in conjunction with the TC-Fe database [18]. In a Scheil simulation, calculations rely on the assumption that no diffusion occurs in the solid and that perfect mixing occurs in the liquid due to efficient stirring, which is adequate for the rapid cooling of a hardfacing product

Figure 3Basic principle of cored wire manufacture (courtesy Welding Alloys Limited).

The wire is then baked, drawn to final size and spooled. The nature and amount of powders inser-ted in the sheath per unit length is determined with the help of a dedicated computer program which enables the undiluted composition to be anticipated. Raw materials are carefully selected to reduce the deposited metal inclusion content and avoid micros-tructural inhomogeneities.

3.3 prODuCtiON OF tESt SpECiMENS

Unlike a large majority of hardfacing products, 167Nb-S remains machinable using carbide-tip or ceramic tools. All specimens were therefore machi-ned to the required dimensions. Strain hardened material was subsequently ground off using silicon carbide paper. Metallographic and X-ray diffraction specimens were polished using diamond paste. Abrasion test specimens were produced with dimen-sions 25 x 45 x 10 mm, which enabled a large num-ber of tests to be carried out for a given material by successively milling abraded surfaces.

3.4 ChEMiCAL ANALySiS AND MiCrOStruCturAL ChArACtErizAtiON

Bulk chemica l analyses were obta ined by optical emission spectrometry. X-ray diffraction



CONJONCturEétudes et recherche

3.6 rOLLEr prESS rOLLEr CLADDiNg

The initial roller cladding operation is carried out in a workshop, using a set-up such as the one shown in Figure 5. Subsequent refurbishment is performed in-situ, which considerably reduces machine down time and overheads. In-situ surfacing operations usually do not require special surface preparation other than cleaning to remove foreign matter. 167Nb-S can be clad directly onto the rollers, unless a different hardfacing material is already present, in which case the remainder of the old hardfacing is removed and a buffer material is applied. Rollers are preheated to 300-350 °C and maintained in that temperature range throughout the cladding operation. Layers are built-up with submerged arc cladding (3.2 mm dia-

meter wire is used in conjunction with a basic flux to minimise the inclusion content of deposited metal). Typical voltage and current settings are 30 V and 450 A. With a deposition rate of 7.5 kg h–1, the arc time required to clad a roller press roller is of the order of two days.

4. rEsuLts of thE invEstigation

4.1 MiCrOStruCturE

Results obtained from XRD show that the d-spa-cings of the predominant phase coincide with those of ferrite (e.g. peaks no. 6, 10 and 16 in

Figure 6). XRD also reveals the presence of retained austenite (e.g. peaks no. 5 and 7). Good agreement was found between experimental data and niobium carbide. Interestingly, two overlapping diffraction peaks were recorded for each set of planes (e.g. peaks no. 25 and 26, 27 and 28). For each set, the relative intensity of the first peak (corresponding to a larger d-spacing) is on average 55 % higher than that of the second peak, as shown in Table 3. These measurements strongly suggest the pres-ence of two types of niobium carbides with slightly different d-spacing. A few low intensity peaks were not labelled in Figure 6, most of which could be correlated with the d-spacings of minor phases present (e.g. the intermetallic compound Fe0.62Cr0.21

Mo0.17) [22].

Figure 4Photograph of the rig used to perform three-body abrasion tests and schematic diagram, with: 1 sand hopper; 2 sand nozzle; 3 rubber-coated wheel; 4 specimen; 5 weight.

Figure 5Schematic illustration of a set-up used for cladding rollers, with: 1 roller; 2 hardfaced surface; 3 cladding gun; 4 heating system; 5 drive motor; 6 control box; 7 power source (courtesy Welding Alloys Limited).




Figure 6Diffractometer plot obtained from 167Nb-S. An identification of individual peaks can be found in Table 3.

Electron microscopy revealed that the microstruc-ture of 167Nb-S consists of large niobium carbides, surrounded by a martensitic matrix which is reinfor-ced by much finer niobium carbides. Some retained austenite can also be found along grain boundaries, together with an intermetallic phase, as shown in Figure 7.The specimen used for microstructural observation in the SEM was extracted from a slowly cooled multi-layer deposit. The relatively coarse microstructure thereby obtained was ideal for microanalysis and the identification of individual phases. In actual cladding applications, cooling rates are higher and 167Nb-S exhibits a much finer microstructure, as shown in Figure 8. The mechanisms leading to microstructu-ral refinement at high cooling rates are discussed in the following section. It is also important to note that the microstructure of 167Nb-S is essentially

equiaxed. The large majority of deposits will exhibit dendritic growth in the heat flow direction, which clearly is not the case in the micrograph shown in Figure 8b (which is a cross-section, perpendicular to the top of the deposit).Large niobium carbides (phase a in Figure 7) are found intragranularly and may form as dendritic crystals (a good example of which is shown in the middle of Figure 8a), provided that cooling rates are reasonably low. The distance between the extremities of such dendrites has been observed to reach 400 µm. Cooling rates obtained when cladding large items such as rollers are sufficiently high to prevent the growth of such crystals, as heat dissipation is efficient; these carbides then form as discrete polygonal particles (Figure 8b). EDS results show that large NbC nucleate on alumina impuri-ties. The solubility of Fe, Cr, Mo and V is quite low;

Figure 7SEM micrograph of 167Nb-S, showing the main phases present, with: a large niobium carbide; b martensitic matrix; c fine niobium carbide; d retained austenite; e intermetallic phase; f grain boundary. Secondary electron image.



CONJONCturEétudes et rechercheTable 3: Comparison between the d-spacings of ferrite, austenite and NbC and experimental data obtained from 167Nb-S [26, 28]. Note: RI is the relative intensity

of diffraction peaks; D-spacing notation: D, id{hkl} where D is the d-spacing in Å, i the diffraction peak’s intensity for a family of planes {hkl}

Peak no. 2Θ, deg RI (%) d{hkl}, Å d{hkl} (α), Å d{hkl} (γ), Åd{hkl} (NbC)

Set no. 1, Åd{hkl} (NbC)

Set no. 2, Å

1 34.778 59.2 2.5774 – – 2.5808, 100d{111} –

2 34.890 26.2 2.5694 – – – 2.5808, 100d{111}

3 40.401 34.8 2.2307 – – 2.2352, 80d{200} –

4 40.513 18.3 2.2248 – – – 2.2352, 80d{200}

5 43.290 24.6 2.0883 – 2.0800, 100d{111} – –

6 44.506 100.6 2.0340 2.0268, 100d{110} - – –

7 50.448 6.2 1.8075 – 1.8000, 80d{200} – –

8 58.495 13.3 1.5766 – – 1.5804, 35d{220} –

9 58.636 7.1 1.5731 – – – 1.5804, 35d{220}

10 64.704 7.0 1.4395 1.4332, 20d{200} – – –

11 69.911 13.8 1.3444 – – 1.3476, 26d{311} –

12 70.135 6.5 1.3407 – – – 1.3476, 26d{311}

13 73.467 5.9 1.2879 – – 1.2903, 9d{222} –

14 73.715 3.9 1.2842 – – – 1.2903, 9d{222}

15 74.250 7.1 1.2762 – 1.2700, 50d{220} – –

16 81.987 11.1 1.1743 1.1702, 30d{211} – – –

17 86.924 1.1 1.1198 – – 1.1175, 5d{400} –

18 87.426 0.7 1.1147 – – – 1.1175, 5d{400}

19 90.133 3.1 1.0881 – 1.0830, 80d{311} – –

20 97.653 7.1 1.0233 – – 1.0253, 7d{331} –

21 98.024 5.2 1.0204 – – – 1.0253, 7d{331}

22 98.886 2.7 1.0139 1.0134, 10d{220} – – –

23 101.166 8.1 0.9971 – - 0.9994, 10d{420} –

24 101.555 4.2 0.9943 – – – 0.9994, 10d{420}

25 115.700 7.3 0.9098 – – 0.9124, 7d{422} –

26 116.154 4.4 0.9075 – – - 0.9124, 7d{422}

27 127.737 2.1 0.8580 – – 0.8602, 6d{511} –

28 128.357 1.2 0.8557 – – – 0.8602, 6d{511}

29 137.242 1.7 0.8272 0.8275, 6d{222} – – –

Figure 8Microstructure of 167Nb-S.Microstructure obtained at a low cooling rates; b high cooling rates. Unetched regions consist of niobium carbides, retained austenite and small amounts of intermetallic phase, while the martensitic matrix is darkly etched. Bright field (BF).

a b



any titanium present is however taken into solu-tion, as shown in Figure 9a. Micro-hardness rea-dings performed on these carbides gave an average of 2540±400 HV0.2, which is in good agreement with published data (Figure 1).Finer niobium carbides (phase c in Figure 7) are found to adopt a plate-like morphology and exhibit higher solubility for chromium and vanadium; only minute amounts of titanium are occasionally found in these carbides. Compositional differences between the two types of niobium carbides are highlighted in Table 4. Leaving aside solute elements present in minor concentration, it is important to note that the niobium to carbon ratio widely differs between the two types of niobium carbides: on an atomic basis, the Nb:C ratio is very close to 1 for large NbC, but reaches only about 0.7 for finer NbC. Although EDS does not enable accurate carbon content measure-ments, the difference in stoichiometry was confir-med by numerous readings and explains why two diffraction peaks were recorded by XRD for each set of planes [23].

As shown previously, regions close to grain bounda-ries consist of retained austenite and an intermetal-lic phase. EDS revealed that, besides having higher carbon content than the matrix, retained austenite also contains slightly greater amounts of Si, Cr and Mo. The intermetallic phase was found to be enriched in Cr, Mo and V. Numerous very fine car-bides were observed along grain boundary regions; despite their small size and the poor resolution of EDS at high magnification, results suggest that these are chromium-rich carbides. The presence of exces-sive amounts of retained austenite and intermetallic phase is undesirable in this type of alloy and the volume fraction of both phases was significantly reduced through alloy development.As shown in later sections, 167Nb-S exhibits opti-mum properties in the as-deposited condition, when the hardness of the matrix is high (micro-hardness readings from regions consisting of a mixture of α’ and plate-like NbC gave 1610±70 HV0.2). Micros-tructural data were nevertheless also acquired fol-lowing heat treatment, which are briefly summarised here. XRD and microscopy revealed that the retai-ned austenite fully transforms after tempering at 600 °C for 14 h (with heating/cooling rates set to 20 °C h–1). Together with the tempering of the mar-tensitic matrix, evidence was found of the precipita-tion of chromium-rich M7C3 and M23C6. Being a stable phase, niobium carbide remains unaffected.

4.2 ALLOy SOLiDiFiCAtiON

Niobium carbide is the first solid phase to become stable in alloys whose composition is within the range given in Table 2. In the case of materials with lower alloying content (e.g. 51Nb-S), the bulk of NbC precipitates together with δ at about 1400 °C.

In such materials, niobium carbides are present as uniformly dispersed plates. An increase in niobium and carbon contents raises both the temperature at which NbC becomes stable and its volume fraction. As shown in Figure 10, NbC precipitation in 167Nb-S is expected to start at 1775 °C and the liquid + niobium carbide phase field (L+NbC(p)) extends over 365 °C.Previous micrographs showed that the bulk of niobium carbides form either as large dendritic crystals (Figure 8a) or discrete polygonal particles (Figure 8b), depending on deposit cooling rate. These carbides were designated primary niobium carbides (NbC(p)) in order to distinguish them from NbC forming at lower temperatures. Figure 11 reveals that the volume fraction of NbC(p) precipitating in the L+NbC(p) phase field (i.e. between 1775 °C and 1410 °C) is very


PhaseLarge (dendritic)

NbCFine (plate-like)

NbC

C, wt% (at%) 10.9 (47.5) 13.3 (51.3)

Si, wt% (at%) … 0.3 (0.5)

Cr, wt% (at%) 0.3 (0.3) 1.7 (1.5)

Mo, wt% (at%) 0.5 (0.3) 1.3 (0.6)

Nb, wt% (at%) 83.9 (47.3) 70.4 (35.2)

Ti, wt% (at%) 2.6 (2.9) 0.1 (0.1)

V, wt% (at%) 0.2 (0.2) 0.7 (0.7)

Fe, wt% (at%) 1.6 (1.5) 12.2 (10.1)

Table 4: Typical composition of niobium carbides, based on standardless EDS analysis

Figure 9 EDS plots obtained from 167Nb-S. Typical EDS plots obtained from a primary niobium carbide and b quasi-eutectic niobium carbide, showing compositional differences. Note: the acquisition time was set to 30 s in both cases.

Figure 10Scheil simulation showing the mole fraction of liquid present in 167Nb-S (current optimised composition) as a function of temperature. No phases were suppressed prior to performing the calculations.



CONJONCturEétudes et recherche

large, as it accounts for 78 mole% of all niobium car-bides forming upon solidification, which is consistent with microstructural observations.Solidification proceeds with the simultaneous forma-tion of δ ferrite and quasi-eutectic niobium carbides (NbC(q–e)) until the alloy reaches 1390 °C. Though the growth directions of NbC(q–e) and neighbouring NbC(p) are identical (e.g. Figure 12), NbC(q–e) adopts a different morphology and its composition is signi-ficantly different to that of NbC(p), as was shown in Table 3. Compositional differences measured by EDS (Table 3) were substantiated by the solidification model. A comparison between the calculated com-positions of NbC at 1600 °C and 1400 °C (which cor-responds to the compositions of NbC(p) and NbC(q–e)

respectively) confirmed the higher solubility of V, Cr and Mo in NbC(q–e), as well as the propensity to record a lower Nb:C ratio in NbC(q–e).

Austenite becomes stable at 1390 °C and few nio-bium carbides are expected to form below that temperature. In the model shown in Figure 10, no intermetallic phase was predicted to be stable at the later stages of solidification. This is partly due to the formation of chromium-rich M7C3 (which leaves less Cr in the remaining liquid), but also because of small changes in the alloy composition, which enabled a significant reduction of the volume fraction of inter-metallic phase.On the basis of these results, the solidification of 167Nb-S can be summarised as follows:

to be reduced from 22% to less than 8%, as deter-mined by image analysis. Solidification modelling was also used to adjust alloy composition in order to reduce the amount of intermetallic phase which was quite substantial, as can be seen in Figure 14. The intermetallic phase observed in 167Nb-S becomes stable only at the later stages of solidifi-cation, which is why it is observed along grain boun-daries. Being very brittle, the presence of interme-tallic phase significantly reduces the tensile yield strength of the alloy, which can result in intergranu-lar fracture during solidification or even in service if tensile stresses operate. Changes in alloy composi-tion enabled the intermetallic phase to be virtually eliminated, which solved sporadic deposit cracking problems. Cracking is however not always solely related to alloy composition. Raw material selection and wire manufacturing practices are of great impor-tance. The use of inadequate raw materials can lead to local microstructural inhomogeneities such as

Figure 12Light micrograph of 167Nb-S showing the distribution and morphology of primary and quasi-eutectic niobium carbides. Note that the directions of growth of NbC(p) dendrite arms and those of plate-shaped NbC(q–e) are identical within a grain. AW, slowly cooled weld. BF, surface polished and heated at 325 °C for 3h.

Figure 11Scheil simulation showing the mole fraction of niobium carbide present in 167Nb-S (current optimised composition) as a function of temperature. No phases were suppressed prior to performing the calculations.

1775 °C 1410 °C L + NbC(p) +

L →→→→ L + NbC(p) →→→→ δ + NbC(q-e) 1390 °C L + NbC(p) + δ 1385 °C L + NbC(p) +

→→→→ + NbC(q-e) + γ →→→→ NbC(q-e) + γ 1235 °C L + NbC(p) + NbC(q-e) < 50 °C NbC(p) + NbC(q-e) →→→→ + γ + M7C3 →→→→ + α’ + M7C3 + γret

Thermodynamic calculations show that M7C3 preci-pitates below about 1235 °C. The precipitation of M7C3 does occur in 167Nb-S, but the volume fraction of these carbides is low with respect to the total volume fraction of carbides.

4.3 ALLOy DEvELOpMENt

Modelling not only provided a good understanding of alloy solidification mechanisms, but also enabled the refinement of alloy composition. Early versions of 167Nb-S contained large amounts of retained auste-nite, as can be seen in Figure 13.The presence of a large volume fraction of metastable retained austenite is undesirable, mostly because it lowers the bulk hardness of the alloy, hence its resis-tance to abrasive wear, as will be shown in a later section. Modelling was used to assess the effect of chemical composition upon the stability of γ as a function of temperature. The resulting changes in alloy composition enabled the volume fraction of γret

Figure 13Microstructure of an early version of 167Nb-S, with: a primary niobium carbide; b martensitic matrix; c quasi-eutectic niobium carbide; d retained austenite; e chromium-rich M23C6; f grain boundary. BF, slowly cooled deposit.

Figure 14Microstructure of a more recent version of 167Nb-S, with: a primary niobium carbide; b martensitic matrix; c quasi-eutectic niobium carbide; d retained austenite; e intermetallic phase; f grain boundary. BF, rapidly cooled deposit.

Figure 15Micrograph showing an example of microstructural inhomogeneity (cluster of polygonal niobium carbide). BF, rapidly cooled deposit.



partially melted particles or, as shown in Figure 15, the formation of NbC clusters. The presence of such defects has been linked to deposit cracking.

4.4 prOpErtiES

As was emphasised in the introduction, high hard-ness is an important requirement for materials to withstand abrasive wear. Besides the effect of composition, the hardness of 167Nb-S deposits was found to be dependent on the cladding proce-dure; variables such as substrate preheat tempera-ture, heat input and the number of layers deposited are of particular importance [24]. The hardness of 167Nb-S normally ranges between 56 and 60 HRC. In the case of lower hardness deposits, measure-ments performed on clad components showed that the hardness can increase to 63 HRC in service as a result of strain hardening. The relationship observed between volume loss and hardness for alloys whose composition is within the range shown in Table 2 is illustrated in Figure 16.

two or three successive layers of metal to obtain the desired properties (sometimes, even a single layer is sufficient). In such instances, the last layer is slightly diluted with the substrate; the hardness of single or two layer deposits is usually somewhat lower than that obtained for multi-layer deposits. This is the case for HC-O, TiC-O and CN-O, which are two-layer deposits. As all three materials are highly susceptible to solidification cracking, abrasion test specimens inevitably exhibited fine cracks. Test specimens manufactured from 51Nb-S and 167Nb-S deposits were crack-free and had an undiluted che-mical composition.The fact that 167Nb-S can be deposited crack-free has many benefits, one of which is related to wear rates. A macroscopic examination of the abraded surfaces of materials which do exhibit cracks clearly showed that cracks and other surface defects such as porosity increase wear rates. Abrasive particles were found to cause accelerated wear at the widest openings of cracks. Conversely, wear was uniform along the abraded surface of 167Nb-S.


Figure 18SEM micrographs of abraded surfaces, with a 167Nb-S (AW condition, slowly cooled), showing a series of protruding dendritic NbC(p) secondary arms (e.g. a), numerous NbC debris (e.g. b) and a NbC trail (c), which illustrates the gradual material removal process large and deep scar obtained with low carbon steel; b a lower alloy material (51Nb-S, tempered at 650 °C for 15 h), showing evidence of carbide spalling (see cavities a and b) presumably caused by a single abrasive particle (see scar c). Secondary electron image.

Figure 16Relationships between adjusted volume loss and hardness for alloys in the Fe-Cr-Nb-C system (undiluted deposits).

Whereas wear rates are low when the hardness of deposits is greater than about 57 HRC, the abra-sive wear resistance quickly decreases below that value. The graph shown in Figure 16 is based on data acquired from testing slowly cooled deposits. For a given hardness, rapidly cooled deposits exhibit even lower wear rates. Two abraded surfaces are shown in Figure 17. Some of the abrasive wear mechanisms operating were uncovered by studying abraded surfaces in the SEM. In the case of 167Nb-S, the SEM revealed a gradual material removal process characterised by the progressive wear and fracture of primary niobium carbides, as shown in Figure 18a. Although primary NbC protrude from the surface, the hard martensi-tic matrix clearly provides the necessary support for these carbides, as no evidence of spalling could be observed. Scars caused by passing silica abrasives were found to be relatively shallow.

Figure 17Macrophotographs of two abrasion test specimens, with a large and deep scar obtained with low carbon steel; b much shallower scar produced when testing 167Nb-S.

Results were also acquired for heat treated speci-mens and can be summarised as follows: a drop in bulk hardness of 5 HRC was recorded after tempe-ring at 550 °C for 5 h. The hardness then remained unchanged for tempering times up to 15 h. Volume losses recorded following heat treatment (at 550 °C for 15 h) were almost twice as high as those obtai-ned with specimens not subjected to heat treatment. Wear mechanisms were found to be very different following heat treatment: an examination of abraded surfaces in the SEM revealed that primary NbC are either extensively damaged by abrasives or knocked out of the matrix (e.g. Figure 18b). The tempered martensitic matrix also sustained much damage, which accounts for greater material loss during the abrasion process.The resistance of various materials to three-body abrasive wear (determined according to ASTM stan-dard G65) is represented graphically in Figure 19. The AWM composition of standard hardfacing mate-rials is shown in Table 1; the nominal composition of 51Nb-S and 167Nb-S is given in the section on alloy design. In practice, many applications require only

Figure 19Compilation of abrasion test results obtained for low carbon steel (LCS) and various hardfacing materials at a given hardness.



conjoncturEétudes et rechercheMachinability is another important property of 167Nb-S which distinguishes this material from conventio-nal hardfacing products. For instance, milling of high carbon/high chromium hardfacing deposits results in extensive tip break off due to high speed impact with large chromium carbides; tools therefore quickly become unusable. Although tool wear was obviously high, 167Nb-S deposits at 60 HRC were machined wit-hout tool break off. In the case of roller-press rollers, this is a clear advantage as run-on and run-off rings present on either side of a refurbished roller can be removed more easily, for instance, by turning.

5. Results obtained in seRvice

An example of roller press roller cladding using 167Nb-S is shown in Figure 20. Data obtained from cement works located in Por-tugal were compiled in Figure 21 in order to show the performance of two rollers in service. An initial cladding operation was carried out on both rollers in 1995 using an early version of 167Nb-S containing a relatively high volume fraction of retained austenite. A refurbishment operation subsequently took place one year later in situ using an improved version of 167Nb-S, in order to restore the roller to its original diameter. Average wear rates were calculated using the following equation:

r = 1 S

i = 1

n

dwi

(3) n dMi

Average wear rates are shown in Figure 21, adjacent to each set of data. In the case of the fixed roller (Figure 21a), wear rates recorded from the initial clad surface were initially of the order of 5.3 mm per million tonnes and with extended service increased to about 20 mm per million tonnes. The roller surface therefore had to be refurbished after 711635 tonnes

fly some components which could benefit from this cladding technology:(i) Cage bars and screw flights in oil presses(ii) Grinder sleeves for clay mills(iii) Solid fuel extruder mouldsExperimental work has already been undertaken with cage bars and very encouraging results have been obtained with an over-alloyed gas-shielded version of 167Nb-S. Crack-free single layer deposits with hardness 61 HRC were thereby obtained, which still exhibited good machinability and grindability.

7. discussion

7.1 Solidification and microStructurE of 167nb-S

As was shown in previous sections, the solidification of 167Nb-S begins at 1775 °C with the heteroge-neous nucleation of primary NbC on alumina impu-rities. All the titanium present goes into solution in NbC(p) due to the low enthalpy of formation of TiC, in addition to the high solubility of Ti in NbC [20, 25]. The kinetics of NbC(p) formation are of particular

Figure 21Graphical representation of wear at the surface of rollers as a function of the mass of clinker ground, with: a fixed roller; b mobile roller (working together). In each graph, wear data obtained from the initial clad surface (early version of 167Nb-S) was superimposed with data obtained from the refurbished surface (optimised version of 167Nb-S). Wear was recorded by the customer using a well established method of measurement; [17] each data point represents the average of 52 measurements on the roller surface.

of clinker had been ground; the wear recorded was w = 7.5 mm. Wear rates recorded after refurbishing the roller with an optimised version of 167Nb-S were approximately 2.7 mm per million tonnes. This large improvement enabled the comminution of over a million tonnes of clinker and the average recorded wear was only w = 2.7 mm. The large reduction in wear rates obtained as a result of material develop-ment are also shown in Figure 21b. The work carried out on 167Nb-S enabled the service life of rollers to be extended considerably, as the amount of clinker ground in the roller press is proportional to the num-ber of hours in service. Though the critical minimum roller diameter had not been reached, subsequent refurbishment took place in situ over a year later, in order once again to restore rollers to their origi-nal diameter. Wear rates obtained thereafter were essentially identical to the ones recorded after the first repair.

6. new applications

Given the unusual combination of properties 167Nb-S has to offer, it is useful to mention brie-

Figure 20Roller press roller clad with 167Nb-S, with: a photograph of a new roller, clad using the set-up shown in Figure 5 (two stringer bead layers, followed by a zig-zag pattern). Roller diameter: 1.2 m; b Close-up of the surface, showing the high quality crack-free deposit obtained.

a b



importance with regard to the final microstructure obtained. In slowly cooled deposits, the nucleation rate of NbC(p) is lower, but subsequent growth is rapid, which is why large NbC crystals form. Conver-sely, the nucleation rate of NbC(p) is high and growth is slow when cooling rates are higher, which results in the formation of a large number of discrete and closely spaced NbC. As mentioned earlier, an exami-nation of deposit cross-sections often reveals den-dritic growth in the heat flow direction, which clearly does not occur in 167Nb-S. This difference is due to the presence of NbC(p), which acts as an inoculant in the melt for the nucleation of δ ferrite when the deposit temperature reaches about 1410 °C. A more isotropic equiaxed microstructure is thereby obtai-ned. A grain refinement effect is also observed with increasing cooling rate, as the number of NbC(p) par-ticles increases.The formation of δ ferrite is accompanied by that of quasi-eutectic NbC between 1410 °C and 1390 °C. The sudden precipitation of NbC(q–e) is believed to occur as a result of the low solubility of carbon in δ ferrite. As δ forms, carbon is rejected at the L/ δ interface, where plate-like NbC(q–e) growth occurs. Previous work carried out on Nb-containing high speed steels showed that the mole fraction of plate-like NbC(q–e) is a function of the mole fraction of δ ferrite forming upon solidification (xδ) [20]. The matrices of alloys in which little or no δ forms upon solidification were found to have a far less uniform carbide distribution. Solidification modelling revea-led that in order to promote these carbides, xδ can be raised most effectively by increasing the Nb content of the alloy and decreasing the carbon, although an increase in other alloying elements such as Si, Mo and V is beneficial.Solidification proceeds with the formation of auste-nite, while the remaining liquid is enriched in carbon, silicon, chromium, molybdenum and vanadium. The nature and volume fraction of phases forming during the last stages of solidification is highly dependent on composition. In the optimised version of 167Nb-S, the enriched liquid gives rise to small amounts of inter-metallic phase and retained austenite. Earlier versions of 167Nb-S exhibited a high volume fraction of inter-metallics, which was due to the segregation of large amounts of Cr and Mo to grain boundaries. Excessi-vely high carbon content in early alloys (with respect to the amount of niobium present) led to a very large volume fraction of retained austenite, while Cr-rich M23C6 precipitated along grain boundaries in alloys containing in excess of 8 wt% chromium.

7.2 rELAtiONShipS MiCrOStruCturES/ prOpErtiES

The presence of relatively coarse primary NbC undoubtedly is beneficial in terms of wear resis-tance. Wear rates recorded for materials containing

very little primary NbC (e.g. 51Nb-S) are compara-tively high. The effectiveness of large primary NbC in 167Nb-S was revealed by electron microscopy. Inevitably, primary NbC particles are much smaller when deposit cooling rates are high; these carbides are however more closely spaced and abrasion tests confirmed that wear rates are even lower than those recorded for slowly cooled deposits, at a given bulk hardness. The presence of quasi-eutectic NbC is also of great importance as these carbides contribute to both strength and wear resistance of the matrix. Electron microscopy showed that the martensitic matrix of 167Nb-S provides the necessary support for carbides. EDS measurements revealed that the matrix is sufficiently alloyed to obtain extensive precipitation of fine carbides upon heat treatment, which may be used to further improve wear resis-tance. Although a slight increase in hardness may be obtained through tempering at moderate tem-perature and short times (i.e. 525 °C < θ < 550 °C, t < 5 h), this was not seen as a viable option, as such a heat treatment is considered impractical with large parts [17]. The presence of excessive amounts of retained austenite and intermetallic phase along grain boundaries were shown to be undesirable. As a general guideline, optimum wear resistance is obtained from deposits without heat-treatment containing a low volume fraction of retained auste-nite and with bulk hardness in excess of 57 HRC. Deposit cracking problems are avoided when the microstructure is uniform and the volume fraction of intermetallic phase is low.

7.3 pErFOrMANCE OF 167NB-S iN SErviCE

In practice, hardfaced surfaces are those of engi-neering components which are often themselves subjected to mechanical loading. In such instances, the structural integrity of the component is of great importance. Most wear-resistant hardfacing pro-ducts exhibit stress relief cracking; the presence of cracks is clearly undesirable as it can lead to crack propagation during service, which eventually induces spalling or even catastrophic failure.Some consumables used in the past to clad roller-press rollers could not be deposited using submer-ged-arc techniques (e.g. TiC-containing materials) due to slag detachability problems; such materials exhibited solidification cracking problems, porosity and high inclusion content. It has been demonstra-ted through numerous refurbishment operations that 167Nb-S can be deposited crack-free, provided that cladding procedures are followed. In addition to this, high quality linings are obtained with 167Nb-S by submerged arc cladding. Component service life benefits from both of these fundamental differences, as uneven wear due to the presence of defects does not occur. In terms of abrasive wear resistance,

results obtained in service are more than satisfac-tory, as confirmed by roller users.For technical reasons, bolted-on segments cannot be substituted for cladding on those roller presses already equipped with hardfaced or studded rol-lers.17 The short term forecast for roller cladding using 167Nb-S is therefore extremely favourable.

8. concLusions

The present paper has summarised theoretical and experimental work carried out to develop a wear resistant hardfacing based on the Fe-Cr-Nb-C system. Microstructural characterisation and solidification modelling enabled an optimisation of the composition of the alloy. Laboratory wear tests and microscopy uncovered essential relationships between micros-tructure and abrasive wear resistance. Owing to its unique set of properties, 167Nb-S is successfully replacing more conventional hardfacing materials for the refurbishment of roller press rollers. Results obtai-ned in service are quite outstanding and the service life of many other components subjected to severe abrasive wear could be extended by adopting this cladding technology, particularly if structural integrity and machinability are sought.

9. rEfErEncEs

[1] “Hardfacing catalogue - cored welding wires”, Welding

Alloys Group, August 2001, 8-9.

[2] A.R. Lansdown and A. L. Price: in “Materials to resist

wear - A guide to their selection and use”, 45-58; 1986,

Oxford, Pergamon Press.

[3] I.M. Hutchings: in “Tribology: friction and wear of engi-

neering materials”, 133-141; 1992, London, Edward

Arnold.

[4] S. Atamert: in “Stability, wear resistance, and micros-

tructure of iron, cobalt and nickel-based hardfacing

alloys” (PhD thesis), The University of Cambridge, 1988,

19-21, 27-31, 60-76.

[5] K.H. Zum Gahr: Metal Progress, September 1979, 116,

(4), 46-52.

[6] H.R. Shetty, T.H. Kosel and N.F. Fiore: Wear, September

1982, 80, (3), 347-376.

[7] J.K. Fulcher, T.H. Kosel and N.F. Fiore: Wear, February

1983, 84, (3), 313-324.

[8] F. Heisterkamp and S.R. Keown: Metals and Materials,

October 1978, 12, 35-36.

[9] S.R. Keown: Symposium “High Speed Steel”, Lång-

shyttan, Sept. 1981, Fagersta, 29-37.

[10] H.J. Goldschmidt: in “Interstitial alloys”, 92-95; 1967,

London, Butterworths & Co.

[11] E.K. Storms: in “The Refractory Carbides”, 11, 12, 55, 56,

71, 112, 134; 1967, London, Academic Press.

[12] A. Kagawa, T. Okamoto, K. Saito and M. Ohta, Journal of

Materials Science, August 1984, 19, (8), 2546-2554.




conjoncturEétudes et recherche[13] K.H. Zum Gahr and D.V. Doane, Metallurgical Trans-

actions A, April 1980, 11A, (4), 613-620.

[14] C.P. Cookson: Surfacing journal, 1985, 16, (4), 125-127.

[15] F. L. Smidth, “HRP. Hydraulic Roller Press”, July 2000.

[16] “Economical comminution with Köppern high-pressure

roller presses”, Köppern, 1998, 2, 6.

[17] C.J. Christensen: Conf. Proc. ”F.L. Smidth International

Conference Seminar”, Copenhagen, 1998, 1, (Ch. 15),

11-14, 26, 27.

[18] Division of Computational Thermodynamics, Department

of Materials Science and Engineering, Royal Institute of

Technology, Stockholm, Sweden.

[19] D.A. Porter and K.E. Easterling: in “Phase transformations

in metals and alloys”, 2nd edn, 208; 1992, London, Chap-

man & Hall.

[20] F. Scandella, “Développement d'un acier rapide pour le

revêtement de cylindres de laminage à chaud”, Soudage

et Techniques Connexes, Institut de Soudure, Yutz, 64,

March/April 2010, 35-46.

[21] “Annual Book of ASTM Standards”, Standard G 65-94,

232-243; 1994, Philadelphia, PA, ASTM.

[22] XRD data: Cr12Fe36Mo10, ref. 31-0401, 1998, JCPDS-Inter-

national Centre for Diffraction Data.

[23] S.J.B. Reed: in “Electron microprobe analysis”, 2nd edn,

274-276; 1993, Cambridge, Cambridge University Press.

[24] F. Scandella: “Results of the experimental program on

the latest version of 167Nb-S”, Welding Alloys internal

report, 1997.

ABSTRACT en fRAnçAiS

Cette publication présente de façon succincte les travaux théoriques et expérimentaux qui ont été menés pour développer un produit de rechargement dur désigné “167Nb-S”. Ce matériau, qui est basé sur le système d'alliage Fe-Cr-Nb-C, a été conçu pour le revêtement de composants soumis à de fortes contraintes d'abrasion en service. Les travaux menés incluent la caractérisation microstructurale, la simulation numérique de solidification et des essais d'abrasion. La microstructure de dépôts réalisés avec le 167Nb-S est constituée d'une fraction volumique importante de carbures de niobium primaires dans une matrice martensitique contenant de fins carbures de niobium quasi-eutectiques. Un peu d'austénite résiduelle peut être présente sur les joints de grains, avec également des composés inter-métalliques en faible proportion. Grâce à des itérations de compositions, il a été possible de réduire la quantité d'austénite résiduelle et de phase intermétallique initialement présente, ce qui se traduit par une amélioration de la résistance à l'abrasion des dépôts et une réduction de leur sensibilité à la fissuration. Des essais normalisés d'abrasion à trois corps ont montré que des dépôts de 167Nb-S ont une résistance à l'usure comparable à celle de dépôts de fonte au chrome hypereutectique ayant un niveau de dureté proche. Le 167Nb-S offre de nombreux avantages par rapport à des produits de rechargement conventionnels tels que la possibilité de réaliser des dépôts multicouches à 61 HRC sans fissuration et ces dépôts restent usinables. L'une des grandes applications de ce produit est le revêtement des cylindres de presses à cylindre, en substitution de produits de rechargement contenant du titane ou encore de segments.

[25] R.W.K. Honeycombe and H.K.D.H. Bhadeshia: in “Steels

microstructure and properties”, 2nd edn, 66; 1995, Lon-

don, Edward Arnold.

[26] XRD data: a-Fe, ref. 06-0696, 1996, JCPDS-International

Centre for Diffraction Data.

[27] XRD data: Austenitic stainless steel, ref. 31-0619, 1996,

JCPDS-International Centre for Diffraction Data.

[28] XRD data: NbC, ref. 38-1364, 1997, JCPDS-International

Centre for Diffraction Data.


études et recherche - · PDF fileand deposit cracking susceptibility. Standard...

Documents

Transcript of études et recherche - · PDF fileand deposit cracking susceptibility. Standard...